_ -:7--:-:;-_/_ __ }L-_:_ : NASA Coh_erence P_a_i;on_056 ........... .......

N

ectrochemical

....................... Resea rch

- .... _ "--_\" L" • ° _ -=---- " -

'..... eCh ogy ............ = .....

....... . .... _.......... : . . " 2Z.........

_ - _..i: :_ _ - _ _._ 7-+!.: ---- ..... :--......=: -.........--

-- - 2__y--_2 __- //- o : ..... - ... "_:" ......... - Proceedings of a conference held at

" . v -- ̧- -_ .... -- - __ ....... :_. : , _ -- --_ = ..... _ • . _,Y'_ _ . _ - •

391 p_ C.gCL lOh NgU-Z0q-RZ

Unc I a s

H]./4e, 0 2 6.-'5I. _4
m

_ ,--=

n

II

iI' I INlI |

_-_ ......... _____ I
NASA Conference Publication 3056

Space

Electrochemical

Research

and

Technology

(SERT)

1989

Richard S. Baldwin, Editor

Lewis Research Center

Cleveland, Ohio

Proceedings of a conference held at

NASA Lewis Research Center

Cleveland, Ohio

April l 1-13, 1989

National Aeronautics and

Space Administration

Office of Management

Scientific and Technical

Information Division

1989
i

is

_z
FOREWORD

Electrochemical energy storage devices will continue to play an important

role in our nation's space program. Space power technology requirements for

future spacecraft and missions are diverse and rapidly expanding; higher

power and energy levels will be required, and the challenges to decrease mass

and to increase efficiency, life, and reliability will remain as significant

goals. The development of advanced electrochemical energy storage

technologies and systems must continue in order to meet NASA's mission needs

and to enable new space opportunities in the future.

A wide variety of advanced electrochemical system concepts are in various

stages of development, new cell electrochemistries are being studied in the

laboratory, and programs which are directed towards the optimization of

components and the development of new materials are actively being pursued.

High energy density secondary batteries and regenerative fuel cell systems

currently under development, as well as other advanced energy storage concepts

which are being conceived and investigated, may represent the enabling

technologies for future space missions and extraterrestrial surface power

applications.

The second NASA Space Electrochemical Research and Technology (SERT)

Conference was held at the Lewis Research Center on April 11-13, 1989. The

objectives of the conference were to examine current technologies, research

efforts, and advanced concepts, and to identify technical barriers which

affect the advancement of electrochemical energy storage systems for space

applications. The conference provided a forum for the exchange and discussion

of ideas and opinions of a proper mix of research scientists, mission

planners, technology advocates, and power system engineers, representing

forty-eight industrial, academic, and governmental organizations.

The conference opened with a series of related overviews in the areas of

NASA advanced mission models (Low-Earth-Orbit (LEO) and Geosynchronous-EarthOrbit (GEO) Missions, Unmanned Planetary Missions, and Human Exploration

Mission Studies). The technical sessions that followed consisted of papers

grouped into the following theme areas:

(I) Advanced Concepts

(2) Hydrogen-Oxygen Fuel Cells and Electrolyzers

(3) Nickel Electrode

(4) Advanced Rechargeable Batteries

The technical sessions were chaired by highly respected investigators who are

extremely knowledgeable in the practical aspects of complete electrochemical

system considerations.

Following the technical presentations, concurrent workshop sessions were

held to discuss in depth the critical issues that related to the technical

theme areas. These workshops, which were led by the technical session

chairpersons, provided a forum for general discussions to coalesce views and

findings into conclusions on progress in the field, prospects for future

advances, areas overlooked, and the directions of future efforts.

PRECEDING PAGE BLANK NOT FILMED
This proceedings includes the related overviews, the technical papers,

and the viewgraphs of the workshop summary presentations, which were prepared

by the session chair- and co-chairpersons.

On behalf of the Power Technology Division and the Lewis Research Center

staff, I would like to thank all of the conference attendees for their active

participation and contributions. Sincere appreciation is also extended to

Dr. Patricia O'Donnell, who chaired the conference, and to the staf_ of the

Electrochemical Technology Branch, Jo Ann Charleston, the Lewis Conference

Planning Office, and ERCI/D-K Associates, Inc. fortheir assistance and

guidance.

Richard S. Baldwin

1989 SERT Conference Coordinator

|V

m_--
CONTENTS

Page

RELATED OVERVIEWS

OVERVIEW OF BATTERY USAGE IN NASA/GSFC LEO AND GEO SPACECRAFTS

Thomas Y. Yi, NASA Goddard Space Flight Center ............

THE IMPORTANCE OF BATTERIES IN UNMANNED MISSIONS

John W. Klein, Jet Propulsion Laboratory ...............

HUMAN EXPLORATION MISSION STUDIES

Robert L. Cataldo, NASA Lewis Research ................

ADVANCED CONCEPTS

A SURVEY OF ADVANCED BATTERY SYSTEMS FOR SPACE APPLICATIONS

Alan I. Attla, Jet Propulsion Laboratory ...............

ADVANCED RECHARGEABLE SODIUM BATTERIES WITH NOVEL CATHODES

S. Di Stefano, B.V. Ratnakumar and C.P. Bankston, Jet Propulslon

Laboratory ..............................

MONOLITHIC SOLID OXIDE FUEL CELL DEVELOPMENT

K.M. Myles and C.C. McPheeters, Argonne National Laboratory .....

THE ELECTROCHEMCIAL GENERATION OF USEFUL CHEMICAL SPECIES FROM LUNAR

MATERIALS

Kan J. Tsai, Daniel J. Kuchynka and Anthony F. Sammells, Eltron

Research, Inc .............................

A NEW CONCEPT FOR HIGH-CYCLE-LIFE LEO: RECHARGEABLE MnO2-HYDROGEN

A.J. Appleby, H.P. Dhar, Y.J. Kim and O.J. Murphy,

Texas A&M Unlversity .........................

CHARACTERIZATION TESTING OF A 40 AHR BIPOLAR NICKEL-HYDROGEN BATTERY

Jeffrey C. Brewer, NASA Marshall Space Flight Center, and

Michelle A. Manzo and Russel P. Gemeiner,

NASA Lewis Research Center ......................

ADVANCED DOUBLE LAYER CAPACITORS

S. Sarangapani, P. Lessner, J. Forchione, A. Griffith and

A.B. LaConti, Giner, Inc .......................

17

33

4]

51

6l

69

83

HYDROGEN-OXYGEN FUEL CELLS AND ELECTROLYZERS

RECENT ADVANCES IN SOLID POLYMER ELECTROLYTE FUEL CELL TECHNOLOGY

WITH LOW PLATINUM LOADING ELECTRODES

Supramaniam Srlnivasan, David 3. Manko, Hermann Koch,

Mohammad A. Enayetullah and A. John Appleby, Texas A&M University 95
THEAPPLICATIONOFDOWCHEMICAL'SPERFLUORINATED MEMBRANESIN

PROTON-EXCHANGE MEMBRANEFUELCELLS

G.A. Eisman, The DowChemical Company ................

HYDROGEN-OXYGEN PROTON-EXCHANGE MEMBRANEFUELCELLSANDELECTROLYZERS

R. Baldwin, NASALewls ResearchCenter, M. Pham,NASALyndon B.

Johnson Space Center, and A. Leonida, J. McElroy and T. Nalette,

Hamilton Standard ..........................

ELECTROCATALYSIS FOROXYGENELECTRODESIN FUELCELLSANDWATER

ELECTROLYZERS FORSPACEAPPLICATIONS

Jal Prakash, Donald Tryk and Ernest Yeager, Case Western Reserve

University ..............................

OXYGENELECTRODESFORRECHARGEABLE ALKALINEFUELCELLS-II

L. Swette and N. Kackiey, Glner, Inc .................

115

127

137

149

NON-NOBLEELECTROCATALYSTS FORALKALINEFUELCELLS

S. Sarangapanl, P. Lessner, M. Manouklanand ,I. Giner,

Glner, Inc .............................. 165

CORROSIONTESTINGOFCANDIDATESFORTHEALKALINEFUELCELLCATHODE

Joseph Singer and William L. Fielder, NASALewis ResearchCenter 171

NICKELELECTRODE

STRUCTURAL COMPARISON OF NICKEL ELECTRODES AND PRECURSOR PHASES

Bahne C. Corniisen, Xiaoyln Shan and Patricla Loyselle,

Mlchigan Technologlca] University .................. 181

IMPEDANCE STUDIES OF Ni/Cd AND Ni/H 2 CELLS USING THE CELL CASE

AS A REFERENCE ELECTRODE

Margaret A. Reid, NASA Lew_s Research Center ............. 197

THE APPLICATION OF ELECTROCHEMICAL IMPEDANCE SPECTROSCOPY FOR

CHARACTERIZING THE DEGRADATION OF Ni(OH)2/N_OOH=ELECTRODES

D.D. Macdonald and B.G. Pound, SRI International, and

S.J. Lenhart, Ford Aerospace ..................... 21l

KOH CONCENTRATION EFFECT ON THE CYCLE LIFE OF NICKEL-HYDROGEN CELLS.

iV. RESULTS OF FAILURE ANALYSES

H.S_ Lim and S.A. Verzwyve]t, Hughes Aircraft Company ........ 243

ADVANCES IN LIGHTWEIGHT NICKEL ELECTRODE TECHNOLOGY

Dwaine Coates, Gary Paul and Paul Daugherty, Eagle-Picher

Industries, Inc ............................ 259

ADVANCED RECHARGEABLE BATTERIES

MULTI-MISSION Ni-H 2 BATTERY CELLS FOR THE_I990's

Lee Miller, Jack Bril] and Gary Dodson, EagIe-Picher

Industries, Inc ............................ 269

vi
NICKEL-HYDROGEN CAPACITYLOSSONSTORAGE

M|chelle A. Manzo, NASALewis Research Center ............ 279

SODIUMSULFUR-BATTERYFLIGHTEXPERIMENTDEFINITIONSTUDY

RebeccaChangand Robert Minck, Ford AerospaceCorporation ..... 297

ADVANCEDHIGH-TEMPERATURE BATTERIES

P.A. Nelson, Argonne National Laboratory ............... 307

ADVANCESIN AMBIENTTEMPERATURE SECONDARYLITHIUMCELLS

S. Subbarao, D.H. Shen, F. Deligiannls, C-K. Huangand

G. Halpert, Jet Propulsion Laboratory ................ 325

TECHNICALWORKSHOPSUMMARYPRESENTATIONS

ADVANCEDCONCEPTSWORKSHOP

Chairperson: John Appleby, Texas A&MUniversity .... _ .......

HYDROGEN-OXYGEN FUEL CELLS AND ELECTROLYZERS WORKSHOP

Chairperson" James Huff, Los Alamos National Laboratory .......

NICKEL ELECTRODE WORKSHOP

Chairperson" Albert Zlmmerman, Aerospace Corporation ........

ADVANCED RECHARGEABLE BATTERIES WORKSHOP

Chalrperson" Gerald Halpert, Jet Propulsion Laboratory ........

339

343

347

35l

vll
m

|

|
N90-20455

OVERVIEW OF BATTERY USAGE IN NASA/GSFC LEO AND GEO

Thomas Y. Yi

NASA Goddard Space Flight Center

Greenbelt, Maryland 20771

SPACECRAFTS

A survey of the batteries used in the LEO and GEO missions at the Goddard Space Flight Center is

described. For each spacecraft, tentative launch date is given, along with relevant battery

parameters including battery size and description, depth-of-discharge, predicted mission

temperature, and life requirement.

INTRODUCTION

The Goddard Space Flight Center (GSFC) has a role to expand human knowledge of the Earth, its

environment, the solar system, and the universe. And as such, the Center has been at the forefront

of space research and exploration through a myriad of spacecrafts, both low Earth orbiting (LEO)

and geosynchronous Earth orbiting (GEO) vehicles. A number of these satellites are listed in Table

1 which shows the relevant battery paramters associated with each spacecraft. All the spacecrafts

listed in the table have baselined the nickel-cadmium batteries. All the cells used in these batteries

are manufactured by Gates Aerospace Batteries (GAB) in Gainesville, Florida.

BATIERY USAGE

In July 1989, Cosmic Background Explorer (COBE) will be launched on a Delta rocket from the

Western Test Range. COBE is an in-house program to study the big-bang theory. COBE was

originally scheduled for a shuttle launch and it was designed with two 50Ah NASA Standard

batteries containing 22 cells per battery. However the launch vehicle was changed to a Delta rocket,

and the whole spacecraft was redesigned to fit this vehicle. The battery design was changed to a

modified 20Ah NASA Standard battery, 18 cells per battery. The battery usage profile is benign.

The mission requirement is for 1 year of battery usage with a predicted mission depth-of-discharge

(DOD) of 0 to 24% over one 63-day eclipse season. The batteries will be subjected to a temperature

range of 16 to 22°C. McDonnel-Douglas Company is the prime for the batteries.

The next GSFC spacecraft to be launched is National Oceanic and Atmospheric Administration

satellite (NOAA-D) in December 1989 for weather observation purposes. This spacecraft is among

a series of NOAA weather satellites which GSFC is cooperating with the National Oceanic and

Atmospheric Administration on. NOAA-D is in a LEO/Polar morning orbit with a predicted

mission DOD of 0 to 16%. There are two 26.5Ah batteries in the spacecraft, 17 cells per battery.

The mission duration is 2 years with predicted mission temperature of 5°C. The next spacecraft,

NOAA-I, is scheduled to be launched in May 1989. NOAA-I will operate in a LEO/Polar afternoon

orbit with 18% DOD, 5°C predicted mission temperature, and 2 year mission duration. It will have

three 26.5Ah batteries, 17 cells per battery. NOAA-J is planned for July 1992 launch, and

NOAA-K, November 1993. The batteries planned for the NOAA-K, L, and M spacecrafts will

utilize the GAB lightweight 50Ah NiCd design with salient characteristics of the NASA Standard

cell. General Electric-Astro, East Windsor is the prime contractor for NOAA satellites.
After NOAA-D, GammaRayObservatory(GRO)will belaunchedfrom STS37in April 1990.

Theprimaryobjectiveof thesatelliteisto studythegammarayradiationphenomenon.The

spacecraftis in aLEOorbit withpredictedmissionDODof 15%at 15°C.Thebatteryrequirement

is for 2.5yearsof operation.TherearesixNASA Standard50Ahbatteriesin GRO,22cellsper

battery. Thesebatteriesarepartof two modularpowersystems(MPS)madeby

McDonnel-DouglasCompany,eachMPScontaining3 batteries.

In July 1990,thefirst of thesecondgenerationGeostationaryOperationalEnvironmentalSatellite

(GOES)will belaunchedfor weatherobservationpurposes.This satellite,GOES-I,isin aGEO

orbit with batteryparametersof 60%DOD,7°Caveragetemperature,and5yearslife requirement.

Therearetwo 12Ahbatteriesin thespacecraft,28cellsperbattery.TheremainingGOEShavethe

samebatteryparameters.GOES-Jis scheduledfor November1991launch,andGOES-K,May

1992launch. Thelaunchdatesfor GOES-LandGOES-Mhavenotbeenfinalized. Ford

AerospaceandCommunicationsCompanyis theprimefor GOES.

ThelastGSFCsatellitescheduledfor launchin 1990is theTrackingandDataRelaySatellite

(TDRS)in December1990.TDRS-Ewill belaunchedto aGEOorbit for communicationpurposes.

Thespacecraftwill containthree40Ahbatteries,24cellsperbattery.Relevantbatteryparameters

includethepredictedmissionDODof 50%,missiontemperatureof 5°C,anddurationof 10years.

Thenextsatellite,TDRS-Fis scheduledfor December1992launch,andTDRS-G,May 1994.The

nextgenerationof TDRSwill probablyutilizeothercell designssuchasadvancedNiCdor NiH2

insteadof theexistingNiCddesign.TRWis theprimefor presentgenerationof TDRS.

InAugust1991,thefirst of theGSFCexplorerplatformsatelliteswill belaunched.Extreme

UltravioletExplorer(EUVE), whichis in aLEO regime,will haveoneMPSwhichcontainsthree

NASA Standard50Ahbatteries.Thebatterieswill besubjectedto apredictedmissionDOD of

15%,temperatureof 15*C,andlife of 3 years.

UpperAtmosphereResearchSatellite(UARS)is plannedto belaunchedfrom STS50in December

1991to studytheEarth'sozonelayerandotherenvironmentalconcerns.Thisspace_a-ft]s ha a 56 °

inclination LEO. Like EUVE, UARS contains one MPS. The predicted mission DOD is 0 to 24%,

temperature, 10 to 16"C, and mission duration, 3 years.

In addition to these missions, the GSFC has a number of Small Explorer satellites which will be

launched from Scout rockets. Battery requirements for these spacecrafts have not yet been

finalized.

R

__-

|

i

I
TABLE 1. - OUTLINE OF GSFC MISSIONS AND ASSOCIATED BATTERY PARAMETERS

MISSION

COBE

NOAA-D

GRO

GOES-I

GOES-J

GOES-K

TDRS-E

TDRS-F

TDRS-G

NOAA-I

EUVE

UARS

LAUNCH

7/89

12/89

4/90

5/90

11/91

5/92

12/90

12/92

5/94

5/91

8/91

12/91

BATTERY

2 20 Ah Batteries

18 CeUs/Battery

2 26.5 Ah Batteries

17 CeUs/Battery

6 50 Ah Batteries

22 Cells/Battery

2 12 Ah Batteries

28 Cells/Battery

3 40 Ah Batteries

24 Cells/Battery

3 26.5 Ah Batteries

17 Cells/Battery

3 50 Ah Batteries

22 Cells/Battery

3 50 Ah Batteries

22 Cells/Battery

ORBIT

LEO/Polar

0-24% DOD, 63-day Eclipse

1 Year Mission

16-22"C

LEO/Polar Morning

0-24% DOD

2 Year Mission

5"C

LEO

15% DOD

2.5 Year Mission

15°C

GEO

60% DOD

5 Year Mission

7°C

GEO

50% DOD

10 Year Mission

5°C

LEO/Polar Afternoon

18% DOD

2 Year Mission

5°C

LEO

15% DOD

3 Year Mission

15°C

LEO, 56 ° Inclination

0-24% DOD

3 Year Mission

10-16°C
_=_
N90-20456

THE IMPORTANCE OF BATTERIES IN UNMANNED MISSIONS

John W. Klein

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, California 91109

Introduc_on

The planetary program has historically used batteries to supply peak power

needs for mission specific applications. Any time that additional power has been

required in order to meet peak power demands or those applications where only

limited amounts of power were required, batteries have always been used. Up

until the mid to late 70's they have performed their task admirably. Recently,

however, we have all become aware of the growing problem of developing reliable

NiCd batteries for long mission and high cycle life applications. Today we will talk

about the role rechargeable batteries will play for future planetary and earth

observing spacecraft.

Historical Uses of Batteries

Planetary

Batteries have always been used whenever there was a need for power in

excess of what a standard primary power source could develop, whether it be from

solar arrays or radioisotope thermoelectric generators. Primary batteries have

been used in probes and other small spacecraft where limited mission life (on the

order of hours) is required. Where peak power needs exceed those available from

a primary source and are cyclic In nature, then secondary or rechargeable batteries have been used.

Ranger

Ranger was the first spacecraft to use rechargeable batteries. It utilized a AgZn battery to augment the power produced from solar arrays. No

battery failures occurred during these set of launches.

Mariner

The Mariner series of spacecraft saw the development of the power

system as we know it today. It started with AgZn batteries for Mariners 1

through 7. On Mariner 8, which launched in May 1971, NiCd batteries were

used for the first time. Due to the careful development and preflight testing

PRECEDING. PAGE BLANK NOT FILMED
of these batteries, no failures occurred in the Mariner 7 and 8 power system, which were identical. Similarly, Mariner 9 and 10 flew NiCd batteries

successfully.

Viking

Following the successes of the Mariner series of spacecraft, it was

decided to develop a lander and an orbiter for Mars observation. In both

cases, due to the success of the NiCd batteries of the Mariner series, NiCds

were used. Even on the lander, which had a much different environment

than had been encountered in space, NiCds were the battery of choice. As

in the Mariner case the NiCds performed as expected with no failures.

i :i-i: = _-_

Leo/Geo

Planetary is not the only mission set where batteries have a historical base.

As we are all aware, they have been used successfully in Leo and Geo orbits. The

list of these missions is long and marked with great successes due to the c_eful

planning, engineering, and test of the spacecraft subsystems. To list just a few of

these include Seasat, the TIROS series, and the whole range of our communications capabilities today.

Recent Experiences ,_ _......

_ I wofii-diike_to spend just a little bit of time _to talk about some of _e _more

recent experiences with NiCds, however, and solicit the support of this group of

engineers and scientists to give to the power system designer that critical tool

that has eluded him, a reliable secondary battery.

:=

Magellan

Let us start with a rather benign mission and see how the project insured

proper battery operation. First of all, the mission profile for Magellan is rather

benign for the batteries. The battery life is projected at 2100 cycles at a depth of

discharge no greater than 40%. This has kept the Magellan batteries well out of

the operational area of concern, high cycle count, coupled with large depths of

discharge.

Gro - --

An Earth Obse_ng spacecraft, however, is the first to stress the batteries

directly in the operational area of concern. Its high number of cycles has been a

major concern to the industry and has warranted careful monitoring of the Gro

project. As all of you are aware, the Gro test cells have degraded prematurely

using the standard Crane accelerated test. This implies that the batteries may not

perform as expected or required for the Gro mission.

Mars Observer

Noting this failure, the Mars Observer project has similar concerns. The

Mars Observer mission profiles calls for almost 9000 cycles, well beyond _the area

of concern. To insure proper operation of the batteries, early c00rd!nated efforts

between JPL and the Mars Observer pfime-contr-a_t0r GESASD have led to a new

design that_b_tti parties feel will meet _e mission requK_mehts[ Unfortunately,

6

Bm

=

i
it is not a design that has either been qualified or flown before. Thus, early in the

program, test cells will be procured and cycled to see what operational concerns

arise. Unfortunately, the project has decided not to continue into mission profile

testing.

TOPEX

TOPEX is an Earth Observing mission with a similar type of battery profile.

It has a cycle life requirements of 20,000 cycles with a maximum depth of discharge of 15%. While the DOD is not as severe as the Mars Observer condition,

the large number of cycles Is a real concern. TOPEX and MO have teamed

together to perform cell testing in order to verify the battery design, separator

material, and lifetime. Until these test are completed in 1991, no battery decision can be made for either TOPEX or Mars Observer on the final design.

Solution

Thus, what power system engineers around the US and the world need is a

rechargeable battery that is reliable, that has been qualified, and that can meet the

strenuous requirements as posed by some of the more advanced scientific missions. Thus, I would like to solicit support from this austere body to continue to

develop batteries that return to the reliability levels seem in the late 70's. To do

this, two steps must be taken.

Rigorous Quality Control

First, a rigorous quality control plan must be developed that oversees the

development of new battery designs and insures their reliability. In addition, this

plan must take into account the new developments that will occur over the comIng years. We must incorporate advances as they become available, but what does

not need to occur is the wholesale change of a battery design before it has been

proven.

The recent battery workshops held on the NiCd problem have taken the

lead in understanding and remedying the problems seen today. These efforts

must not stop there but continue to improve the reliability and quality of batteries.

Rigorous Acceptance Testing

In addition to quality control, rigorous acceptance and accelerated life testing must continue on all cells that will fly on strenuous missions. This Is a hard

pill to swallow in today's world where every dollar has to be justified. However, I

hope that the recent GRO experience will convince project managers across the

US to take seriously the advantages of rigorous acceptance and life testing.

Mission Profile Testing

A caveat to the acceptance test concept is to also perform mission

profile testing so that the true nature of the long term performance is

understood. Without performing such long term tests, there is no way to
predict the true operating characteristics of the battery that is in flight. In

addition, If something occurs within the battery and/or the power system

and one battery is lost, having a battery that has gone through some mission

profile testing is invaluable to the understanding of future operational issues.

Of particular interest is how far can the battery be discharged safely, when

does it need to be reconditioned, how can the most energy be extracted

from the battery under the new operating condition.

Closing

i-n conclusion, NiCds have been and will continue to be the mainstay of the

power system engineers tools for peak power production. Recent experience has

tarnished its once sterling reputation. However, the industry has Stood up to this

challenge and implemented wide ranging plans to rectify the situation. These

efforts should be applauded and supported as new designs and materials become

available.

In addition, project managers must become aware of their responsibility to

test "their" batteries and insure quality and mission operating characteristics.

Without this teamwork, the role of NiCds in the future will diminish, and other

batteries, not as optimum for high performance application (low mass and volume)

will take their place. Let us not let this happen, but to continue to strive to

develop batteries that are reliable, rugged, and qualified.
N90-20457

HUMAN EXPLORATION MISSION STUDIES

Robert L. Cataldo

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio 44135

The nation's efforts "to expand human presence and activity beyond Earth

orbit Into the solar system" was given renewed emphasis In January of 1988

when the Presldential Directive on National Space Policy was slgned Into

effect. The expansion of human presence into the solar system has particular

significance, In that It defines long-range goals for NASA's future missions.

To embark and achieve such ambitious ventures is a significant undertaking,

particularly compared to past space actlvltles.

INTRODUCTION

Two major efforts recently released, the National Commission on Space

Report, "Pioneering the Space Frontier," 1986, and astronaut Dr. Sally Ride's

task force report, "Leadership and America's Future in Space," 1987, have

helped set goals and glve focus to NASA's future direction. Dr. Ride's task

force formulated a plan to achieve the Commission's proposed agenda for the

civilian space program. The task force recommended four initiatives: (1) Misslon Planet Earth, (2) Exploratlon of the Solar System, (3) Outpost on the

Moon, and (4) Humans to Mars. The first two initiatives, to a great degree,

are being pursued by NASA's Office of Space Science and Applications. However,

the last two initiatives prompted the task force to recommend that a special

NASA office be established to coordinate and lead human exploratlon studies.

Therefore, the Office of Exploration was established to provide a focal point

for these activities. These Include establishing a mature understandlng of

mission options and opportunltles and deflnlng those near-term activltles that

can provide the greatest impact on future mlsslons.

The Office of Exploration has established a process whereby aIl NASA

field centers and other NASA Headquarters offices participate In the formulatlon and analysis of a wlde range of misslon strategles. These strategies

were manifested into specific scenarios or candidate case studies. The case

studies provided a systematic approach into analyzing each mission element.

First, each case study must address several major themes and rationale IncludIng: national pride and international prestige, advancement of scientific

knowIedge, a catalyst for technology, economlc benefits, space enterprise,

international cooperatlon, and education and excellence. Second, the set of

candidate case studies are formulated to encompass the technology requirement

limits in the life sciences, launch capabllities, space transfer, automation

and robotics, In-space operations, power, and propulsion.

The first set of reference case studies identlfy three major strategies:

(l) human expeditions, (2) science outposts, and (3) evolutionary expansion.

During the past year, four case studies were examined to explore these strategies. The expedltlonary mlsslons Include the Human Expedition to Phobos and

Human Expedltlon to Mars case studies. The Lunar Observatory and Lunar
Outpost to Early Mars Evolution case studies examined the later two strategies. Thls set of case studies establlshed the framework to perform detailed

mission analysls and system engineering to define a host of concepts and

requirements for various space systems and advanced technologles. This paper

describes the details of each mission and, specifically, the results

affecting the advanced technologles required to accompllsh each mission

scenario.

MISSIONSTUDIES

The Office of Exploration has established a process whereby reference

case studies are examined yearly. This allows for a reasonable subset of _ _

options to be studied in some detail in order to establish first order effects

and trends In various system elements. Once the study has been characterized

in sufficient depth, different system elements can be replaced with new

options so that the overall system performance sensitivities can be analyzed,

The broad scope in case studies is engineered to provide_a wlde range of mlsslon options and to avoid making priorities between explorlng the Moon or Mars.

The fo|lowing candidate case studies are the first set of options studied

In 1988. These options are not the final set, In fact, during the coming

years these cases may be refined and new cases added, to form a comprehensive

set of options so that quallfled engineering decisions can be made in the

final selection process for those studies that will be carried forward into a

program phase A/B.

HUMAN EXPEDITION TO PHOBOS

A manned expedition to Phobos, the inner Martian moon, Would establish an

early U.S. exploration leadership. Activities such as exploratlon, resource

survey and science station set up would be conducted. Also, a fleet of small

teleoperated "low-tech" rovers could provide enhanced exploration of the Martian surface. The expedition to Phobos satisfies two objectives, that of

human exploration and prevlously studied robotic missions with sample return

options.

Past spectrai:analysls of both Phobos and Delmos suggest that materials

are present that could Justify leveraglng rocket propellants for future return

trips from Mars. Propel lant production plants could allow vehicles to carry

fuel only for the trans Earth-Mars leg, and refuel at Phobos or Deimos for the

trans Mars-Earth leg. Earth to orbit (ETO) mass could be significantly

reduced, particularly with an evolutionary Mars mission. The Phobos mission

also relleves the pressure on the launch transportation systems, because of

reduced propellant mass required by not landing on Mars itself.

Mission Scenario

The Phobos expedition employs a "spllt/sprlnt" transportation scheme

where a cargo vehicle carrying science equipment, crew return propellant and

Mars rovers would be launched on a minimum energy trajectory. About lO months

later, the vehicle carrying a crew of four on a high-energy sprint class

...._?! '!"i '

lO

Z

i

E

i
trajectory Is launched. Nine months later the crew vehicle mates wlth the

cargo vehlcle In Mars orbit. Two crew members then board an excursion vehicle

and conduct the mission objectives on Phobos. The two remaining crew members

teleoperate rovers on the Martian surface and transfer propellant from the cargo vehlcle. After about 30 days In the Mars system the four crew members return to Earth having been away from Earth a total of 14 months.

HUMAN EXPEDITION TO MARS

Since the invention of the telescope, Mars, with its observed global dust

storms and "canaIi" has been a favorite destination for planetary exploration

and colonization. Mars has been called Earth's sister planet because of their

similarities, although it Is not a very hospitable planet for humans. Scientists have speculated that Mars had more of Earth's characteristics, denser

atmosphere water for example, which could have supported lifeforms early In

Its evolution. Thus, there is a large international segment of the population

that suggests we return to explore Mars with manned mlsslons.

Mlsslon Scenarlo

The Mars Expedition calls for three separate segments launched on successlve opportunities. The first segment Includes a cargo vehicle preceding

the eight member crew on a spiit/sprlnt approach as in the Phobos mission.

However, the cargo vehicle that transports rovers, surface habitats, and exploratlon and science equipment Is placed In Mars orbit. After spacecraft rendezvous, four crew members descend to the surface In the landing craft for a

20 day stay. The exploration party will perform science experiments and possibly explore the surface for water, minerals, volatiles, and biological traces

of present and past llfeforms. The crew will live In the pre-landed habitats

and utilize the rover for extended exploratlon. The remalning crew members In

orbit will transfer propellant from the cargo vehicle, perform orbital science

and support surface operations. The surface crew w111 unite with the orbit

craft and return to Earth. The total length of the mission would span about

14 months. On ensuing trips exploration of Phobos and Delmos would also be

included In the scenarlo along with different landing sites on Mars.

LUNAR OBSERVATORY

The Lunar Observatory, comprised of optical telescope arrays, stellar monitoring telescopes, and radio telescopes would encompass both the radio and

optical spectrum. An observatory, constructed on the far or backside of the

moon, could provide scientists with informatlon orders of magnitude greater

than could be achieved from one in earth or In low-Earth-orbit. Slgnificant

benefits of lunar locations are; circumventing Earth's atmospheric Impediments, shielding of electromagnetic waves and providing a stable foundation

for large instruments and arrays.

Mission Scenario

An estimated four flights to the Moon's far side would be required to

emplace the observatory. Two launches a year, one cargo and one crew, make up

11
the mlssion set. One combined crew/cargo flight per year would follow. The

crew of four would remain on the Moon's surface ooly during the Lunar day

using their lander vehicle as a habitat, similar to the Apollo mlssions. A

permanent base is not required because of the short stay periods for maintaining the observatory about every 3 years and other exploration missions would

land at different sites.

Unpressurlzed rovers would be used as crew transports during the observatory set up and maintenance, as well as automated machinery used for material

handling and scientific equipment transport.

LUNAROUTPOSTTO EARLYMARSEVOLUTION

An evolutionary approach to space exploration is by building on the experlence of living and working on the Moon. The Moonis very attractive to _

accompl|sh this since the transfer time back to Earth IS only several days in

case an emergency situation would occur. A sustained humahpressence on the

lunar surface would provide valuable data on working in reduced gravity, human

factors considerations with long duratlon planetary mlsslons and learning how

to extract and utilize local resources.

A recommendationof the National Commissionon Space (NCOS)was that a

"bridge betweenworlds" be built with the first step to develop the Moonthen

continue on to Mars. This case study deflnes the methodology necessary to

achieve that goal, by building a space exploration infrastructure that would

lead to nearly self-sufflclent space habitation. This, the most complex and

ambitious scenario, stretches the llmits of technological advances in space

constructlon, utilization of Indlgenous resources, life support, artificial

gravity, chemical and electric propulslon, aerocapture, material and propellant

productlon and space power.

Mission Scenario

A series of both piloted and cargo flights would embark for the Moon.

Cargo would be sent on longer transfer trajectories uslng electric propulsion

and crews sent on fast chemically propelled transfer vehicles. It could take

several years to fu11y construct a surface facllity, possibly with expediated

construction by vehlcles operated from Earth. The principle goals of the base

is to produce materials and liquid oxygen used for Mars flights. The effort

to produce these products is leverage against the massotherwise launched from

Earth.

Sometimelater, possibly 6 to lO years, the Mars flights would begin once

capabllitles are sufficient for the successful Mars trip. Flrst, an electric

cargo vehicle would transport surface equipment, communications orbiters,

rovers, Mars-Phobostransfer vehicles, sclentiflc experiments and a Phobos

propellant plant. As the vehicle approaches Mars, it would deploy;

areosynchronous communication orbiters, robotic equipment to Deimos; and a

fuel processing plant to Phobos. The plant would produce fuel for future crew

return flights.

12
During the next Earth-Mars launch opportunity, the lunar electrlc cargo

vehicle will be reused to transfer the crew's transport vehlcle to lunar orbit

for lunar derived oxygen fueling. The cargo/crew vehicle returns to Earth

orbit, and separates. The cargo vehicle stays In orbit for use as the next

ferry trlp to the Moon, and the crew transport Is boarded by the Mars crew.

After a systems checkout, they begin their B month journey. Several options

exist ks to Mars stay times: 60 days, I year, and 2 years. Continued flights

to Mars could occur every Earth-Mars opportunity which is about

26 months apart.

MISSIONTECHNOLOGY REQUIREMENTS

As the case studles becomemoremature and focused, better deflned

requirements will emerge. However, the preliminary analysis of the 1988 case

studies have surfaced many technology areas needing further development.

These major areas include the following: surface exploration systems, fluid

management,propulsion, automation and robotics In-space operations, aerocapture and power. Table I illustates key technologies as evaluated in relation

to the four case studies. Programs such as Pathfinder and the Civil Space

Technology Initiative are planned to advance certain technologies that will

enable these missions. The requirements refined over the next few years will

provide these programs with mission technology goals. On the other hand, the

technologists must predict the degree to which the objective can be met

considering the techological barriers and funding levels so that the proper

estimates can be Integrated into the case studies. Significant scenario augmentation reflecting the change In technology status maybe necessary to

accomplish mission objectives.

POWERSYSTEMREQUIREMENTS

By deslgn, the case studies span a range of complexity with the more

basic being the Expedition to Phobosand the more complex the Lunar Outpost to

Early Mars. The Lewis ResearchCenter has been chartered by OEXPto perform

the power system analyses. These tasks Include doing broad trades as well as

In depth investlgatlons. Major study areas Include power systems for exploration elements such as; electric cargo vehicles, planetary orbit transportation/

staging nodes, mobile equipment resource processing plants, surface science

outposts and habitated bases.

The power levels presented in table II are flrst cut estimates. As future

systems analysis are performed, the values will becomemore mature. Oneobjective of the power system studies is to discover those technologies that provide

the greatest leverage; either by possessing significant massor reliability

advantages, or having a high degree of commonality. Thus, we will be able to

focus our resources on those technologies that have the highest pay off.

13
CRITICAL TECHNOLOGY

CRYOGENIC FLUID MANAGEMENT

AUTOMATED RENDEZVOUS AND DOCKING

AUTONOMOUS ROVERS

MARS AND EARTH AEROCAPTURE

ON-ORBIT ASSEMBLY & CONSTRUCTION

SURFACE EXTRA VEHICULAR ACTIVITY SUITS

SURFACE POWER (INCLUDING SP-IO0)

ADVANCED CHEMICAL PROPULSION

NUCLEAR ELECTRIC PROPULSION

IN-SITU PROPELLANT PRODUCTION

ADVANCED LIFE SUPPORT

LEGENd. O INDICATES REQUIRED TECHNOLOGY

O INDICATES CRITICAL PACING ELEMENT

HUMAN

EXPEDITION

TO PHOBOS

O

O

O

HUMAN

EXPEDITION

TO MARS

O

O

O

O

LUNAR

OBSERV.

O

O

O

LUNAR

OUTPOST TO

EARLY MARS

EVOLUTION

O

0 0

0 0

0

0

0

TABLE I - ASSESSMENT OF PREREQUISITE TECHNOLOGY

14

_=

|
SYSTEM ELEMENT

PLANETARY UTILITY VEHICLE

PLANETARY MINING VEHICLE

PLANETARY CONSTRUCTION VEHICLE

PLANETARY MANNED EXCURSION

LUNAR OBSERVATORY

LUNAR PROCESSING

LUNAR BASE

MARS OUTPOST

MARS BASE

PHOBO_DEIMOS PROCESSING

CARGO VEHICLE

STAGING NODES

FUEL DEPOT

POWER

1 -2kW

5 - 10 kW

15 kW

30 kW

30 - 50 kW

150 - 300 kW

100 kW

20 kW

40 - 50 kW

1000 kW

5000 kW

TIBD

TBD

COMMENT

RECHARGEABLE

RECHARGEABLEJPEAKING

RECHARGEABLE/PEAKING

CONTINUOUS

FULL NIGHT POWER

ELECTRIC/THERMAL

CREW OF EIGHT

40 DAYS

2 YEARS

TABLE II - ESTIMATED EXTRATERRESTRIAL POWER REQUIREMENTS

OF SIGNIFICANT SYSTEM ELEMENTS

]5
El

|

m_

m

R_
N90-20458

A SURVEY OF ADVANCED BATTERY SYSTEMS FOR SPACE APPLICATIONS

Alan I. Attia

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, California 91109

The results of a survey on advanced secondary battery

systems for space applications are presented. Fifty-five

battery experts from government, industry and

universities participated in the survey by providing

their opinions on the use of several battery types for

six space missions, and their predictions of likely

technological advances that would impact the development

of these batteries. The results of the survey predict

that only four battery types are likely to exceed a

specific energy of 150 Wh/kg and meet the safety and

reliability requirements for space applications within

the next 15 years.

1.0 INTRODUCTION

The Jet Propulsion Laboratory, under the NASA Headquarters

sponsorship of the Advanced Battery Concepts Task, recently completed

an evaluation of various advanced battery concepts to replace the

current Ni-H 2 and Ni-Cd space qualified batteries. The goals were: i) to

identify advanced battery systems capable of outperforming present day

batteries by a significant margin; 2) to obtain an accurate estimate of

the anticipated improvements afforded by some technologies; and 3) to

obtain a consensus as to which of the large number of possible systems

are likely to yield the desired improvements with the highest likelihood

of success by the year 2005, if properly funded.

2.0 APPROACH

Following an initial analysis by JPL of various electrochemical

energy storage devices, the opinion of battery experts was solicited

through a 5-part questionnaire. A brief description of each battery

system considered by JPL was included with the questionnaire as

background information, together with estimates of theoretical and

practical energy densities derived from our initial analysis.

Participants were asked to submit comments and answer the questions only

within their areas of expertise.

PRECEDING PAGE BLANK NOT FILMED

]7
3.0 APPLICATIONS REQUIREMENTS

A long shelf life, from 3 to 7 years, is a firm requirement in

most space applications; however, capacity, cycle life and rate

requirements tend to vary depending on the specific application. The

energy storage requirements of six anticipated space missions are shown

in Table i. The requirements vary widely within the following limits:

(a) charge time of 2 hours to 22 days; (b) discharge time of 0.6 hour

to 17 days; (c) cycle life from 80 to 50,000 cycles; and peak power from

0.5 kW to I00 kW.

3.1 Currently Available and Near-Term Systems

The systems currently in use for space applications include the

Ni-Cd, Ni-H 2 and Ag-Zn batteries. Near-termadvanced systems include the

advanced Ni-Cd battery and the bipolar or common pressure vessel Ni-H 2

batteries. JPL estimates of the characterlstics of the space qualified

and the near-term advanced systems are summarized in Table 2. As is

immediately noticed, substantial reductions in the overall weight of the

battery system would result from a battery capable of a specific energy

in excess of 200 Wh/kg. However, the importance of long cycle life,

safety, and reliability cannot be overemphasized; and high energy

density alone cannot be the only factor to be taken into consideration

when assessing the potential of a specific technology for space

applications.

4.0 RESULTS OF OPINION SURVEY

4.1 Respondents Profile

The questionnaire £ogether with background information wassent

to 205 specialists selected from all sec£ors of industry, government,

the universities and research institutes. Fifty-five respondents

returned the questionnaire, including a low percentage of responses

from universities. Table 3 shows a summary of the make-up of the

respondents group.

4.2 Energy Density Critique

As a starting point for the survey, we had identified a total 23

advanced power sources capable of significant improvements over present

day technology. Six types of power systems were represented: aqueous_

molten salt, solid electrolytes, lithium-halogens, lithium-intercalation

systems, and regenerative fuel cells. The participants• were asked to

evaluate the accuracy of our estimates of achievable specific energies

for the proposed systems and to comment in general on the various

systems presented to them for consideration.

Their responses are summarized in Table 4. On average, the

respondents' estimates were slightly more conservative than JPL's

estimates, but the range of estimates is much wider for systems still

in their early stage of development. For example, the respondents

]8
estimate for the advanced nickel-hydrogen system is 75 ± 6 Wh/kg vs

JPL's estimate of 76 Wh/kg. For the solid electrolyte Li/S system the

respondent's estimate is 335 ± 179 compared to JPl's value of 500

Wh/kg.

4.3 Risk to Develop Successful Aerospace Batteries

The respondents were asked to give their best estimates of the

likelihood of developing the proposed battery systems by the years 1995,

2000 and 2005 and to list the main obstacles to be overcome for each

system. The estimates of these probabilities are summarized in Table 5.

The systems with an acceptable risk for development are marked with an

asterisk. Although the ranges of the estimates are fairly wide, certain

trends are clearly evident.

The standard deviations for the estimates are lower for systems

under active and well funded development; the likelihood of their

development by the year 2005 is also high. Regenerative fuel cells,

Advanced Nickel-Hydrogen and Sodium-Sulfur are typical examples and are

all rated high.

Solid electrolyte systems based on Beta" Alumina and metal

chloride cathodes are rated somewhat lower than Na/S but with acceptable

development risks. The same is true for LiAI/FeS2, lithium-intercalation

systems (including those with polymer electrolytes,) and solid

electrolyte fuel cells. Here again the intermediate rating seems to

reflect the lower degree of funding for those systems.

A few systems are consistently rated "high risk", the lithiumhalogens, the solid electrolyte Na/Cl 2 and Li/S, and the molten salt

Be/NiF 2.

A list of the most frequently mentioned comments and perceived

main obstacles on the most promising candidates is presented in Table

6. Li/S and Li/halogens are included in the list although these two

systems were judged poor prospects for development.

4.4 Worthwhile Systems Omitted from JPL List

The respondents were also asked to identify other advanced

battery candidates omitted or overlooked by JPL, and to give their best

estimate of the realizable specific energy for each system.

Within five of the six categories of systems identified as

potential candidates in our questionnaire, the following additional

systems were suggested as having potential for achieving specific

energies approaching 200 Wh/kg:

(a) Molten Salts:

At 240 Wh/kg, the bipolar "Upper Plateau" LiAI/FeS 2 battery

offers outstanding peak power, in excess of 500 W/kg throughout

its discharge period, over a wide range of states of charge.

19
(b) Solid Electrolytes:

Li/O 2 and Ca/O2 are being explored in conjunction with a solid

oxide ionic conductor operating at high temperatures (>700 C).

Practical energy densities in excess of 200 Wh/kg are

conceivable.

(c) Lithium-Halogens:

Systems of the type Lithium/SO 2 inorganic electrolyte/Metal

Halides are considered by some as safer and more practical

alternatives to the lithium-halogens systems. These systems are

capable of achieving 200 Wh/kg at high rates of discharge, but

safety is a concern.

(d) Li/Intercalation Cathodes:

Several systems with metal oxide Cathodes, not mentioned in our

original list of potential candidates, were considered capable

of achieving high specific energy. Li/CoO 2 (150 Wh/kg) is one

such candidate but requires a suitable electrolyte to achieve

long cycle life. Li/MnO_ (currently at 125 Wh/kg) is another and

could reach a substantlally greater specific energy in bigger

cells, due to the Improved packaging factor in large cells. The

cell voltage of 2.8 V for Li/MnO 2 is inside the electrochemical

window of the current most promislng organic electrolytes.

As a whole the class of Lithium/Intercalation Cathodes offers

specific energies from I00 to 200 Wh/kg, and presents

opportunities for both solid polymer electrolytes as well as

ambient temperature conventionai electrolytes.

(e) Regenerative Fuel Cells:

.... The_solid polymer electr01y£e (SPE) H2/O 2 fuel cell is a major

candidate for the space station. A breadboard system successfully

operated for more than i000 cycles at NASA/JSC. NASA is currently

funding a study to conduct a flight experiment on a reversible

regenerative SPE fuel cell with all passive fluid and thermal

controls. For large systems a specific energy in excess of 200

Wh/kg appears quite feasible. In many respects this system is

very similar to the alkaline RFC and could be considered as a

direct replacement for it.

(f) Supercapacitors:

A most interesting development which may impact energy storage

technology in the future is the supercapacitor. To date devices

with specific energies approaching 5Wh/kg havebeen successfully

demonstrated. These devices offer the additional advantages of

ruggedness, high power density (up to 200 W/cc) and potentially

unlimited cycle life.

20
4.5 Suitability of Systems for Space Applications

The fourth question set required the survey participants to

estimate the degree of suitability of the proposed systems for

several space applications. The panelists were requested to rate

each battery system as highly suitable (H), moderately suitable

(M) or not suitable (L) for each of six types of space missions.

The responses are tabulated in Table 7 for each of the most

promising systems. Based on those estimates, the development

priorities are shown in Table 8.

For the six missions listed, the Beta" solid electrolyte (BASE)

systems, Na/BASE/S and Na/BASE/metal chlorides battery are

considered best for four missions; the H2/O 2 alkaline RFC is

ranked best for two missions. The U.P. LiAI/FeS 2 system was found

worthwhile in 3 missions, whereas the Li/TiS2, as a

representative of a lithium/intercalation cathode system, was

deemed useful in three applications. The Ni/H 2 was also highly

rated for nearly all applications but was not considered an

advanced battery candidate because of its lower specific energy

and advanced stage of development. Some of the matches are quite

evident, as discussed below.

(a) Lithium-Intercalation Batteries for the Planetary Rover:

The planetary rover has requirements that are well suited to

ambient temperature lithium-intercalation batteries, both in

terms of cycle life and rate. The fact that these batteries have

no special temperature control requirements, and can be packaged

and temperature controlled like the rest of the vehicle equipment

is an important plus. The fact that they do not require a close

temperature control is an advantage also. Several chemistries are

available, including polymeric electrolytes, giving flexibility

in a final choice.

(b) Regenerative Fuel Cell for Lunar Base:

The lunar mission is unique in that it requires very long

operating times (days vs hours). In such applications the

dominant weight of the energy storage system is in the reactants.

The regenerative H./O 2 fuel cell has a very high specific energy

for thls appl_catlon, approaching 500 Wh/kg, as most of its

weight would be in the light weight reactants and their required

tankage.

(c) Upper Plateau (U.P.) LiAI/FeS 2 for GEO:

The basis for recommending this system for development is its

high energy density capability and its very high expectations of

success. Recent results obtained at Argonne National Laboratory

have shown that the system is capable of a cycle life in excess

of i000 cycles and a specific energy approaching 200 Wh/kg. The

only other candidate system for GEO, aside from the Ni/H 2 system,

2]
would be the Na/S system. Although significantly more dollars

were spent on Na/S, the U.P. LiAI/FeS 2 system shows expectations

of having a greater energy density, lower risk of catastrophic

failure, less risk of premature shorting, better high rate and

peak power capability, less temperature variation during cycling,

and less safety concerns for manned Shuttle launches. In

addition, the U.P. LiAI/FeS 2 system can be activated prior to

launch.

(d) Na/BASE/S or FeCl 2 or NiCI 2 for LEO and Planetary Orbiters:

The requirements for LEO are very strenuous with respect to cycle

life and can only be met by very few systems (Ni-Cd or Ni-H2!.

Currently the Na/S system is being developed for LEO but is

still considered a high risk due to its high operating

temperature. The Na/Metal Chlorides are attractive alternatives

for this application because of their lower operating temperature

and lower current density, hence lower risk of premature failure.

4.6 Technological Breakthroughs

Finally an estimate of the likelihood of occurrence of several

new technological breakthroughs by the years 1995,.2000 and 2005

was requested from the participants. Oniy those estimates which

show a reasonable degree of certainty are included in Table 9.

a)

b)

c)

d)

e)

5.0 CONCLUSIONS AND FUTURE WORK

The results appear to support the following conclusions:

Most experts believe that the technical problems of the sodiumsolid electrolyte systems, if continued to be funded at their

current level, will be resolved by 2005. This type of system will

provide an intermediate specific energy storage device (130

Wh/kg).

Key requirements for the de_elopment_0f_lithium-intercalation

systems appear well underway t0be_lngresolved by 2005, the main

obstacle at this time being the plating efficiency of the lithium

electrode.

The molten salt U.P. LiAI/FeS 2 system will be a serious high

specific energy (180 Wh/kg) contender by 2005.

The development of a passive _regenerative fuel cell is very

likely by 2005. However, the parallel development of a

bifunctional oxygen electrode is far from being certain.

The halogen-based systems, whether with lithium or sodium, are

unlikely to be developed by the year 2005.

22
f) In general, the survey results are conservative. Participants in

the survey are pessimistic about the chances of success of high

risk developments with high potential payoffs, such as the

lithium/solid ion conductor/sulfur or lithium-halogen battery,

maybe due to the current low level of funding for their

development.

As a follow up to this survey, the participants will be given a

chance to comment on the results and conclusions.

23
TABLE 1

ENERGY STORAGE REQUIREMENTS OF SIX ANTICIPATED

SPACE MISSIONS

PRIORITY CHARGE/

DISCHARGE

DURATIONS

TYPICAL

OPERATIONAL

CYCLES_QUIRED

ACTUAL QUAL* DESIRED

TYPICAL

PEAK POWER AND

ENERGY STORAGE

REQUIRED

#i Outer C - 2 hr 500 1,000 2,000

Planetary D - 0.7 hr

Orbit

#2 Inner C - 2 hr 3,000 6,000 I0,000

Planetary D - 0.7 hr

Orbit

#3 GEO C - 22.8 hr 1,000 1,500 4,000

D - 1.2 hr

#4 Planetary C - 12 hr 300 600 800

Rover D - 3 hr

#5 Lunar C - II Days 80 160 350

Base D - 17 Days

#6 LEO C - 1 hr 30,000 35,000 50,000

D - 0.6 hr

0.5 c (1 KWH)

1.5 C (2 KWH)

1.5 c (5 KWH)

1.3 c (3 KWH)

0.02 c (5 MWH)

i.i C (25 KWH)

GEO = Geosynchronous Orbit

LEO = Low Earth Orbit

*QUAL = Minimum number of cycles needed to qualify for application

=

24
TABLE 2

SECONDARY BATTERIES FOR SPACE APPLICATIONS

SYSTEMS CURRENTLY IN USE

SYSTEM SPECIFIC

ENERGY

(Wh/kg)

Ni-Cd 34*

Ni-H 2 (IPV) $ 45

Ag-Zn 90

ENERGY

DENSITY

(Wh/1 )

7O

25

8O

CYCLE

LIFE

(40% DOD)

m_w..uumm

20,000**

25,000

5O

OPERATING

TEMPERATURE

(°c)

I0 - 20

i0 - 20

I0 - 20

ADVANCES IN STATE-0F-THE-ART

Advanced Ni-Cd 36 ii0

Ni-H 2 (CPV) # 60 70

35,000

15,000

i0 - 20

i0 - 20

* NASA standard, 50 Ah battery

** Standard Cell Qualification Test

$ IPV = Individual Pressure Vessel

# CPV = Common Pressure Vessel

25
TABLE 3

PROFILE OF CONTRIBUTORS TO SURVEY

ON ADVANCED BATTERIES

Occupation

Aerospace Industry

Battery Manufacturers

Other Manufacturers

Government

UniversitiesInstitutes

Number of Respondents

6

15

8

14

3

9

TOTAL 55

26

=m

_=
TABLE 4

ESTIMATES OF ACHIEVABLE SPECIFIC ENERGY FOR ADVANCED BATTERIES

(Wh/kg)

System JPL Panel's Number of

Estimate Estimate Range Responses

AgO/Fe

Advanced Ni/H z

90 85 +/-13 60-110 18

76 75 +/-6 60-80 25

U.P.LiAI-FeS z

Be-NiF z

LiAI-NiS 2

180 154 +/-33 75-188 31

185 156 +/-43 95-185 9

180 155 +/-35 75-184 19

Na/BASE/S

Na/BASE/CI 2

Na/BASE/TCNE

Na/BASE/CuCI z

Na/BASE/FeCI 2

Na/BASE/NiCI 2

Li/Solid Ion

Conductor/S

Lithium/Polymer

electrolyte

130 132 +/-26

200 197 +/-70

i00 95 +/-12

160 132 +/-30

150 130 +/-20

160 137 +/-24

500 335 +/-179 100-510

250 183 +/-67 50-250

80-220 29

100-350 9

70-100 i0

80-160 ii

80-150 19

80-160 17

14

28

Li/CI 2

Li/Br 2

500 375 +/-173 80-500

200 170 +/-56 70-250

15

14

Li/TiS z

Li/NbSe 3

Li/Mo6S 8

Li/V205

Li/a-Cr308

90 95 +/-13 73-130

i00 105 +/-15 80-150

140 126 +/-28 50-180

150 143 +/-27 75-200

200 176 +/-37 75-200

28

25

22

24

22

Zn/O 2 i00

Alkaline RFC i00

Solid Oxide HJO 2 200

99 +/-15 60-140

152 +/-113 100-500

252 +/-180 120-750

17

14

i0

BASE = Beta" alumina solid electrolyte

27
TABLE 5

LIKELIHOOD OF DEVELOPMENT OF ADVANCED BATTERY SYSTEMS

System

AgO/Fe

Advanced Ni/H2*

U.P.LiAI-FeS2*

Be-NiF 2

LiAI-NiS 2

Na/BASE/S*

Na/BASE/CI 2

Na/BASE/TCNE

Na/BASE/CuCI2*

Na/BASE/FeCl2*

Na/BASE/NiCI2*

Li/Solid Ion

Conductor/S

Lithium/Polymer

electrolyte*

Li/CI 2

Li/Br 2

Li/TiS2*

Li/NbSe3*

Li/M06S 8

Li/V2Os*

Li/a-Cr308*

Zn/O 2

Alkaline RFC*

Solid Oxide HJO2*

Estimates of Probability of Development by

1995 2000 2005

37 +/- 32

56 +/- 25

47 +/- 33

71 +/- 19

54 +/- 38

81 +/- 17

36 +/- 26

12 +/" 16

28 +/- 25

52 +/- 31

25 +/- 24

41 +/- 28

60 +/- 30

35 +/- s5 -

46 +/- 31

56 +/- 28

17 +/- 18

23 +/- 20

27 +/- 19

36 +/- 23

35 +/- 21

72 +/- 25

3l +/- 25

33 +/- 28

47 +/- 26

53 +/- 28

54 +/- 25

80 +/- 22

29 +/- 29

42 +/- 33

60 +/- 29

66 +/-29_

66 +/- 27

18 +/i 19 29 +/- 25 39 +/' 28

33 +/- 29 45 +/- 31 55 +/- 30

i0 +/- 14

12 +/- 13

19 +/- 18

21 +/- 18

32 +/- 25

33 +/- 26

44 +/- 26

42 +/- 27

34 +/- 28

34 +/- 23

27 +/- 18

57 +/- 29

57 +/- 31

50 +/- 3z

48 +/- 28

43 +/- 24

66 +/- 31

66 +/- 30

58 +/- 34

58 +/- 32

55 +/- 28

33 +/- 23

49 +/- 30

30 +/- 24

44 +/- 25

63 +/- 28

47 +/- 27

58 +/- 24

77 +/- 26

59 +/- 30

* These sys£e_s present acceptable development risks

28

i

=

i

=

=

m

_=

--_=

m

Z

=

i

E
TABLE 6

OBSTACLES TOWARD THE DEVELOPMENT OF ADVANCED BATTERIES

System Obstacle

Advanced Ni/H2: 2-electron nickel electrode is unlikely

No obstacle to improved Ni/H2 (CPV and/or Bipolar)

U.P. LiAI FeS2:

Na/Beta"/S or Metal

Chlorides

Corrosion and materials compatibility

Reliability of ceramic electrolyte and seals

Lithium-Sulfur:

Thin Film lithium/

polymer electrolyte:

Lithium-Halogens:

Lithium-Intercalation

Cathodes:

Alkaline RFC:

Solid Oxide Fuel

Cell:

Development of lithium ion conducting electrolyte

Need for higher ionic conductivity in polymer

Low cycle life

Material compatibility and corrosion

Lithium cyclability

Poor electrolyte stability

Development of the oxygen bifunctional electrode

Materials compatibility

29
TABLE 7

SUITABILITY OF ADVANCEDBATTERY SYSTEMSFOR SPACEAPPLICATIONS

(L= Not Suitable; M= Moderately Suitable; H= Highly Suitable)

SYSTEM

PLANETARY

INNER OUTER GEO ROVER LUNAR LEO

ORBIT ORBIT BASE

Advanced

Ni/H2

U.P.LiAI-FeS 2

Na/BASE/S

Na/BASE/FeCI 2

Na/BASE/NiCI 2

Lithium/Polymer

electrolyte

Li/TiS 2

Li/V205

Li/a-Cr308

Alkaline RFC

Solid Oxide

H2/O2

L=5 L=7 L=2 L=I0 L=I4 L=3

M=9 M=5 M=6 M=9 M=8 M=7

H=I3 H=I4 }{=20 H=9 H=7 H=20

L=7 L=7 L=8 L=8 L=8 L=I8

M=I3 M=II M=7 M=5 M=I2 M=7

H=6 H=8 H=II H=I4 H=6 H=2

L=5 L=3 L=2 L=3 L=5 L=I3

M=I5 M=I5 M=I3 M=i6 L=II L=i2

H=IO H=I2 H=I5 H=II H=I6 H=6

L=5 L=3 L=2 L=4 L=6 L=I0

M=II M=I2 M=I2 M=I3 M=8 M=8

H=6 H=6 H=6 H=4 H=9 H=4

L=5 L=3 L=2 L=3 L=6 L=9

M=I0 M=I0 M=i0 M=II M=7 M=8

H=7 H=8 H=8 H=6 H=9 H=5

L=18 L=11 L=10 L=9 L=12 L=23

M=5 M=9 M=I0 M=i0 M=6 M=I

H=4 H=6 H=6 H=8 H=9 H=3

L=21 L=II L=II L=I0 L=I6 L=28

M=8 M=I5 M=I3 M=I3 M=8 M=5

H=4 H=7 H=8 H=I0 H=6 H=I

L=24 L=I5 L=I3 L=I2 L=I6 L=30

M=8 M=I4 M=I2 M=II M=II M=5

H=I H=3 H=8 H=9 H=6 H=0

L=21 L=I2 L=II L=9 L=I3 L=26

M=8 M=I4 M=I4 M=II M=II M=5

H=I H=4 H=5 H=9 H=6 H=0

L=9 L=8 L=6 L=6 L=I L=9

M=II M=I2 M=I3 M=I6 M=8 M=9

H=3 H=3 H=5 H=2 H=I6 H=6

L=I0 L=9 L=7 L=8 L=I L=9

M=7 M=7 M=8 M=9 M=5 M=7

H=3 H=4 H=6 H=4 H=I7 H=6

30

z

i

=

=

=
TABLE 8

RANKING OF ADVANCED BATTERY SYSTEMS

Mission Recommended Systems (Ranking)

Planetary

Inner Orbit

Planetary

Outer Orbit

GEO

Planetary

Rover

Lunar Base

LEO

Na/BASE/S or FeCI 2 or NiCI 2 (i)

U.P. LiAI-FeS 2 (2)

H2/O 2 Alkaline RFC (3)

Na/BASE/S or FeCI 2 or NiCI 2 (I)

U.P. LiAI/FeS 2 (2)

Li/TiS 2 (3)

Na/BASE/S or FeCl 2 or NiCI 2 (i)

H2/O 2 Alkaline RFC or U. P. LiAI/FeS 2 (2)

Li/TiS2 (3)

Na/BASE/S or FeCI 2 or NiCl 2 (i)

U.P. LiAI/FeS 2 (2)

Li/TiS 2 or Alkaline RFC (3)

Hz/O 2 Alkaline RFC or Solid Oxide (i)

Na/BASE/S or FeCl 2 or NiCl 2 (2)

U.P. LiAI/FeS 2 (3)

H2/O 2 Alkaline RFC or Solid Oxide (I)

Na/BASE/S or NiCl 2 or FeCl 2 (2)

The advanced Ni-H 2 system was not included in the rankings as this system

is much further ahead in its development stage.

3]
TABLE 9

PROJECTEDp_:0sPECTs FOR_TECHNOLOGICALBREAKTHROUGHS_

% Probability

Future Event 1995 2000 2005

Improved Beta" alumina suitable

for prismatic cells

Improved ceramic seals for Beta" Alumina

Hermetic seals for high temperature

molten salt lithium batteries

A metal oxide intercalation electrode for

lithium batteries with cycle life of i00

A reversible lithium electrode capable of

i000 cycles in organic electrolyte at 25C

(e.g., Li/TiS 2 or Li/Metal Oxide)

a) Medium voltage cells ('2 V.)

b) High voltage cells (-3V.)

Chemical overcharge protection for

Li-AI/FeS 2 molten salt system

The development of very thin, suitable

polymer electrolytes for thin Li cells

Development of hot-launch, ready to use,

high temperature rechargeable batteries

a) Na/Beta" Alumina/S

b) Li-AI/FeS 2

Development of practical rechargeable

designs for

a) Li/CI 2

b) Li/Br 2

High rate (!00mA/cm2), reversible, long

life, oxygen electrode for alkaline fuel

cells

39+/-26

57+/-27

29+/-32

44+/-26

42+/-28

33+/-24

46+/-26

41+/-29

50+/-29

44+/-30

11+/-13

16+/-17

36+/-21

53+/-27

71+/-27

61+/-31

60+/-25

57+/-28

47+/-25

59+/-25

52+/-28

59+/-29

54+/-29

24+/-18

28+/-20

48+/-27

66+/-27

82+/-24

71+/-28

73+/-30

66+/-29

59+/-28

69+/-26

66+/-28

70+/-29

69+/-31

i

33÷/-22

37+/-23

61+/-29

32

i
N90-20459

ADVANCED RECHARGEABLE SODIUM BATTERIES WITH NOVEL CATHODES

S. Di Stefano, B.V. Ratnakumar and C.P. Bankston

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, California 91109

Abstract

Various high energy density rechargeable batteries are being

considered for future space applications. Of these, the sodium

sulfur battery is one of the leading candidates. The primary

advantage is the high energy density (760 Wh/kg theoretical).

Energy densities in excess of 180 Wh/kg have been realized in

practical batteries. More recently, cathodes other than sulfur are

being evaluated. We at JPL are evaluating various new cathode

materials for use in high energy density sodium batteries for

advanced space applications. Our approach is to carry out basic

electrochemical studies of these materials in a sodium cell

configuration in order to understand their fundamental behaviors.

Thus far our studies have focused on alternate metal chlorides such

as CuCl 2 and organic cathode materials such as TCNE.

Introduction

Rechargeable batteries play a significant role in NASA's space

programs including planetary missions and satellite applications such

as Geosynchronous Earth Orbit (GEO) and Low Earth Orbit (LEO)

missions. There is an increasing need for a light weight high energy

battery with long active and cycle lives. Typically, the efforts at

various organizations are aimed at developing a battery system with

an energy density _ 200 Wh/kg at power densities _ i00 W/kg, cycle

life of i000 cycles and active life of ten years. The interest for

many years has been focussed on sodium-sulfur battery which has a

high formula energy density (-750 wh/kg) a good portion of which has

been realized in practical cells for several hundreds of

charge/discharge cycles (I). The battery typically operates at 300-

400°C where the 8" alumina solid electrolyte (BASE) is highly

conductive to sodium ions allowing high power densities drawn from

the battery. Together with these advantages, there are certain

deterrent factors associated with the use of the sodium sulfur

battery which partly protracted the development. R. M. Dell et

al. (2) have given an excellent account of the status of the NaS

battery in their recent review. Many of these operational and

safety problems stem from sulfur electrode, e.g., the excessively

corrosive nature of the discharge product, sodium polysulfide,

towards steels, the high vapor pressure of liquid sulfur above 700°C,

the possible degradation of BASE in polysulfide melts, (3) the

insulating nature of liquid sulfur, in particular the way in which it

33
isolates regions of BASE during recharge, etc. Alternatives to

liquid sulfur cathode are therefore being sought and there are some

examples in the literature based on fused salt electrolytes (4-6).

In this paper, we review our studies on alternate organic and

inorganic cathodes in sodium batteries. In both cases, the cell

contains sodium tetrachloroaluminate as the secondary molten

electrolyte and BASE as primary electrolyte.

Carbonitriles are a class of relatively stable organic compounds

that have high reduction potentials and energy densities. Tetracyanoethylene (TCNE) in particular has an energy density of -620

Wh/kg at a high potential of -3.0 V vs. Na+/Na and a low melting

point. Therefore, TCNE has been examined as a possible liquid

cathode in sodium batteries. Among the inorganic cathode materialS,

transition metal dichlorides (e.g., FeCl 2 and NiCl2) are gaining

increasing attention as solid, insoluble rechargeabie cathodes in

place of liquid sulfur (8-11). These systems can operate at lower

temperatures -170°C, where the corrosion problems are proportionately

less severe. The safety problem is considerably reduced due to the

inherent sluggish kinetics of the chemical reaction between liquid

sodium and NaAICl 4 molten electrolyte in the event of ceramic failure

(12). Further, the molten electrolyte remains invariant during

charge/ discharge preventing polarization losses and localized high

current densities on the BASE surface. High energy densities -150

Wh/kg and long cycle life have been demonstrated _ with ferrous

chloride and more recently with nickel chl0rides. We summarlze below

our electrochemical studies on the ferrous chloride and nickel

chloride cathodes and also on another promising metal chloride,

copper chloride.

Experim_n_a ! .

.... The experimental cell idopted for th_s_-;;{Ehdles is described in

detail in. (7) The cell essentially has a central-cathode design

with a provision to seal the anode half-cell. .Thg_cell is heated

electrically with a hea_ng £ape wound around a stainless steel tube

acting as a container for sodium. For the studies involving TCNE,

either Pt or C is used as the positive current Collector. Metal

chloride positive electrodes were initially fabricated by

impregnating commercially available sintered nickel grfds (used in

Ni-Cd batteries) with the depolarizer by immersing the grid into a

concentrated chloride solution. Copper electrodes cannot however be

made in this manner for long-term studies due to the preferential

dissolution of nickel to copper in the impregnating solution.

Perforated copper grids were therefore used to make electrodes in the

discharged state either by hot pressing with a binder (Teflon) or by

sintering the powders in an inert atmosphere at temperatures below

the melting point of NaCl. Conventional electrochemical equipment

consisting of E G & G 273 potentiostat/Galvanostat and E G & G 5301

lock-in amplifier interfaced with an Apple IIe computer were used.

AC and dc measurements were carried out using E G & G ac impedance

and Softcorr software. Various techniques such as potentiodynamic

polarization, dc linear polarization and ac impedance were employed

to obtain information on the kinetic parameters of the cathode

34
reactions. Discharge/charge cycling was also carried out

galvanostatically in a controlled fashion using the above

electrochemical equipment. All the chemicals TCNE, NaAICI4, NaCl and

metal powder were of analytical grade and were used as received.

BASE tubes were cleaned by etching in hot phosphoric acid. All the

cell fabrication operations were carried out in an argon-filled

glove box.

Results and Discussions

The findings of our studies on TCNE as a cathode material in

sodium batteries are described in detail elsewhere. (7) Briefly,

TCNE reduces at 3.0 V vs. Na+/Na reversibly with a coulombic yield of

>60% at 1 ma/cm 2 on either the Pt or C current collector. Fig. 1

gives the typical voltage-time curves of TCNE during discharge and

charging. The exchange current density on either Pt or C is of the

order of 10 -3 A/cm 2 in a catholyte containing typically 10% of TCNE.

The kinetics of TCNE reduction are essentially governed by the rate

of mass transfer in the cathode as evidenced from ac impedance and dc

potentiodynamic polarization data. In TCNE-rich catholytes there is

a kinetic hindrance to the reduction of TCNE reflected in increased

polarization (ohmic, concentration as well as charge transfer)

losses. As a result, the coulombic yield as well as diffusion

coefficient of TCNE decrease with increasing TCNE concentration in

the catholyte. The sluggish kinetics at higher TCNE concentrations

is related to the degradation processes occurring in TCNE at

operating temperatures of -230°C forming an adduct. FTIR studies

point to a possible polymerization of TCNE under the experimental

conditions. Due to the sluggish kinetics at higher TCNE

concentrations, it may not be possible to achieve high power

densities from TCNE. On the other hand, for medium to low power

applications, TCNE is a good cathode material, since the electrochemical activity and reversibility of TCNE are retained even after

polymerization.

Metal Dichloride Cathodes

Figure 2 shows the typical charge/discharge curves of a "16 mAh

NiCl 2 electrode (made b_ impregnating NiCl 2 into sintered Ni plaque)

at 200°C and at 2 mA/cm z in the first two test cycles. The discharge

as well as charging curves are flat at nominal voltages of 2.56 and

2.63 V vs. Na+/Na respectively. There is not much polarization loss

from the open circuit voltage of 2.59 V.

The coulombic yield is nearly 100% and there is no loss in

capacity during cycling in the first few test cycles. Figure 3

illustrates the high rate discharge capability of the electrode. The

discharge curves at higher current densities, i.e., 4 mA/cm 2 and 8

mA/cm 2 are also flat with mid-point voltages of 2.54 and 2.5 V and

coulombic yields of 98% and 90% respectively. Ferrous chloride

electrodes (made by sintering dry powders of sodium chloride and

iron) exhibit a lower operating potential of -2.3 V as may be seen

from the discharge and charge curves of a "230 mAh electrode in

Figure 4. The discharge curve is flat in this case also, with a

35
coulombic yield in excess of 90%. The copper chloride cathode is

expected to give higher voltage than nickel chloride, from the

electrode potentials of Cu+2/Cu and Ni2+/Ni couples. Accordingly,

higher energy densities may be possible with copper chloride as

compared to ferrous chloride or nickel chloride especially at high

power densities due to a high electronic conductivity of copper

formed in the course of discharge. The essential requirement for a

successful operation of metal chloride cathode in the above sodium

batteries is total insolubility of the chloride in a NaAICI 4 melt.

Analytical studies carried out on the catholyte with CuCl 2 and NiCl 2

revealed no extra dissolution of copper over nickel. Initial studies

with copper chloride (electrode made by impregnation of CuCl 2 into

sintered Ni plaque) show a higher operating voltage than NiCl 2 or

FeCl 2 (Figure 5). The discharge (as well as charging) curve is,

however, not flat in this case as with FeCl 2 or NiCl 2 which may be

due to the mechanism of reduction being different with the possible

formation of a monovalent copper.

DC polarization studies were carried out on these cathode

materials to obtain information on the kinetics of their reduction.

The potentiodynamic polarization curves of NiCl2, FeCl 2 and CuCl 2

(Figure 6) illustrate the high rate capability of these electrodes.

There exist mass transfer effects at high discharge rates interfering

with the charge transfer processes due to which limiting currents are

observed at high overpotentials _ 300 mY. The limiting currents may

be due to a relatively slow diffusion processes inside the porous

cathode. Nevertheless, the limiting current densities are reasonably

high and are of the order of 15-25 mA/cm 2. The charge transfer

kinetics appear to be rapid as evidenced by high exchange current

densities of 2 mA/cm 2, 0.65 mA/cm 2 and 0.4 mA/cm z for NiCl2, FeCl 2

and CuCl2, respectively, based on the geometric areas of these

electrodes. Linear polarization experiments (i.e., at low

perturbations) were also carried out to derive similar information on

£he exchange current densities. The current potential curves at

these potentials are linear (e.g., FeCl 2 in Figure 7) the slopes of

which give the polarization resistances. The exchange current

densities thus obtained are 4.6 mA/cm 2, 1.5 mA/cm 2 and 0.39 mA/cm 2

for NiCl2, FeCl 2 and CuCl2, respectively.

The _complex piane impedance spectrum of NiCI 2 (Figure 8) shows a

large inductance at high frequencies not shown in the figure

attributed to the porous nickel grid and a low frequency semi-circle

ascribed to charge transfer kinetics. The exchange current density

calculated therefrom is 6.9 mA/cm 2 . There is no evidence of

diffusional impedance in the frequency range studied (i.e., _ 0.01

Hz). The complex plane impedance spectrum of CuCl 2 (sintered)

electrode, on the other hand, contains two relaxation loops. The

exchange current density evaluated from the first relaxation loop is

0.9 mA/cm 2. it is however not clear yet as to whether the second

relaxation loop is related to the formation, if any, of monovalent

copper during the reduction of CuCl 2.

36
Conclusions

Among the cathode materials studied at JPL as alternatives to

sulfur cathode in sodium batteries, TCNE is suitable for low to

medium power applications. Metal dichloride cathodes, on the other

hand can deliver energy densities comparable to sodium-sulfur

batteries at high power densities without many of the operational and

safety problems associated with the use of sodium-sulfur batteries

and are thus suitable for various applications including space, EV

load-levelling, etc. The electrochemical studies carried out so far

indicate high exchange current densities for the reduction of these

cathodes.

Acknowledqement

The work described here was carried out at the Jet Propulsion

Laboratory, California Institute of Technology, under contract with

the National Aeronautics and Space Administration. One of the

authors (B. V. Ratnakumar) acknowledges the National Research Council

for providing his Research Associateship during this work.

References

i. J. Sudworth and A. R. Tilley (Eds.), "The Sodium/Sulfur Battery,"

Chapman and Hall, New York 1985.

2. R. M. Dell and R. J. Bones, B. Electrochem., 4, 319 (1988).

3. M. Liu and L. C. De Jonghe, J. Electrochem. Soc., 135, 741

(1988) .

4. J. Werth, U.S. Patent 3,877,984 (1975).

5. G. Mamantov, R. Marassi, M. Matsunaga, Y. Ogata, J. P. Wiaux and

E. J. Frazer, J. Electrochem. Soc., 127, 2319 (1980).

6. K. M. Abraham, L. Pitts and R. Schiff, J. Electrochem. Soc., 127,

2595 (1980).

7. B. V. Ratnakumar, S. Di Stefano, G. Nagasubramanian, R. M.

Williams and C. P. Bankston, Extd Abstr. #74, ECS Fall Meeting,

Chicago, IL, October 9-14, 1988. Manuscript submitted for

publication.

8. J. Coetzer and M. Thackeray, U.S. Patent 4,288,506.

9. J. Coetzer, J. Power Sources, 18, 377 (1986).

i0. R. C. Galloway, J. Electrochem. Soc., 134, 286 (1987).

ii. B. J. Jones, J. Coetzer, R. C. Galloway and D. A. Teagle, J.

Electrochem. Soc., 134, 2379 (1987).

12. R. T. Wedlake, A. R. Tilley and D. A. Teagle, B. Electrochem., _,

41 (1988).

37
z

z

4000- _ '1 _ / I

3560 - 3

. 2680 "_ "_1"_\

U,,I

224 "_"-_.

\\

180t _ I I __..t .....

0 6 12 18 24 30

CAPACITY, mAh

Figure I. Typical Charge/Dlscharge Curves of TCNE

(i0 W%) at 230"C: (I) Fourth blschargs,

(2) Fifth Discharge, and (3) Charge

3000

28O0

z 2600

+

Z

N"

E

2400

LLI

2200

I I I I

2000 __L___.___I I

0 30 60 90

1, min

Figure 3.

120 15£

Discharge Curves of 16 mAh Nicl 2 Electrode

at 200"C and at Current Density of

(I) 2 mA/cm 2 , (2) 4 mA/cm 2 , and (3) 8 mA/cm 2

300O

2800 -

+

Z

2400-

LU

2200

2000.-

0

Figure 2.

I I I I

3

1 I I I

30 60 90 120 150

t, rain

Typical Charge/Discharge Curve 16 mAh

Electrode at 200"C During (i) Second,

(2) Third, (3} Fourth Discharge, and

(4) Charge at 2 mA/cm 2

Z

Z

O9

>

UJ

O

>

3.0

2,8

2.4

2.0

INITIAL CHARGE (FORMATION)j/

FIRST DISCHARGE_

AC AND DC EXPERIMENTS \

AT DIFF POTENTIALS a,

1__ I I I I 1._ I

30 60 90 120 150 180 210 240

CAPACITY, mAh

Figure 4. Initial Charge/Discharge Curves of

FeCI 2 Cathode

i

E

38

-z
3100

2860

2620

Z

E 2380

w

2140

1900

Figure 5.

I t' I I

| . _L __L.. L

0 80 160 240 320

t, min

Charge/Discharge Curves of CuCI 2 Electrode

(Made by Impregnating CuCl 2 into a Sintered

Nickel Grid) at 200"C

28001

2680

2560

+

_ 2440

g

W

2320

2200

........ I ........ I ........ I ........ I ........ I .......

2080

100 101 102 103 104 105

I (i,A/cm 2)

Figure 6. Potentiodynami_: Polarization Curves of

(i) CuC]2, {2) NICl2, and (3) FeCI 2

1oG

U.I

2384

2376

2368

2360

2352 --

2344

-6000

Figure 7.

I I I I

I I I I

-3000 0 3000 6000

I (pA/cm 2)

Linear Polarization Curve for FeCI 2

from which an Exchan@e Current

Density of 1.5 mA/cm z is Calculated

06

E

= I

(a)

0.8

",,.:.:.

e_

Z' (ohms)

I I

2

E

6O

50 (b)

40

30

2O

tO ooooeoeoo

_f.l.____._ I _ I

10 20 30 40 50 60

Z' (ohms)

Figure 8. Complex Plane _mpedance ot (a) NlCl 2

and (b) CuCI 2

39
Z

=z

II

--_i
N90 - 20460

MONOLITHIC SOLID OXIDE FUEL CELL DEVELOPMENT

K.M. Myles and C.C. McPheeters

Argonne National Laboratory

Argonne, Illinois 60439

INTRODUCTION

The Monolithic Solid Oxide Fuel Cell (MSOFC) is being developed by a recently

formed team consisting of Argonne National Laboratory, Allied-Signal

Aerospace/AiResearch, and Combustion Engineering. The MSOFC (ref. i) is an oxideceramic structure in which appropriate electronic and ionic conductors are

fabricated in a "honeycomb" shape similar to a block of corrugated paperboard. The

electrolyte that conducts oxygen ions from the air side to the fuel side is yttriastabilized zirconia (YSZ). All the other materials are electronic conductors,

including the nickel-YSZ anode, the strontium-doped lanthanum manganite cathode, and

the doped lanthanum chromite interconnect (bipolar plate). These electronic and

ionic conductors are arranged to provide short conduction paths to minimize

resistive losses. The power density achievable with the MSOFC is expected to be

about 8 kW/kg or 4 kW/L (ref. 2) at fuel efficiencies over 50%, because of small

cell size and low resistive losses in the materials. The MSOFC operates in the

range of 700 to IO00°C, at which temperatures rapid reform of hydrocarbon fuels is

expected within the nickel-YSZ fuel channels.

Tape casting (ref. 3) and hot roll calendering (ref. 4) are used to fabricate

the HSOFC structure. The tape casting process consists of spreading a ceramic

slurry (slip) uniformly on a substrate, such as glass or polymer film, using a

doctor blade. After the slip is dry, the ceramic layer is stripped off the

substrate in the form of a flexible "tape." The hot roll calendering process

consists of mixing ceramic powder with organic binder and plasticizer and rolling

the warm mixture into a thin tape. The green tapes are cut to the desired

dimensions, and the electrode tapes are corrugated to form the gas flow channels.

The corrugations are formed by folding the tape onto a warm mold. After cooling,

the tape retains the corrugated shape. Corrugated electrode tapes and flat

electrolyte and interconnect composite tapes are stacked up to form the MSOFC

structure. The layers are bonded together in the green state by heating the polymer

binder slightly above its glass transition temperature under a small weight. The

bonded green structure is then heated according to a precise firing schedule to the

sintering temperature to form the monolithic ceramic structure.

The performance of the MSOFC has improved significantly during the course of

development. The limitation of this system, based on materials resistance alone

without interfacial resistances, is 0.05 ohm-cm 2 area-specific resistance (ASR).

The current typical performance of MSOFC single cells is characterized by ASRs of

about 0.4 to 0.5 ohm-cm 2. With further development, the ASR is expected to be

reduced below 0.I ohm-cm z, which will result in power levels greater than 2.5 W/cm 2.

PRECEDING PAGE BLANK NOT FILMED

4]
SYSTEM PERFORMANCE

Because of the high temperature of operation of the MSOFC and the compact fuel

cell design, systems that incorporate the MSOFC for power production are very

simple. The high temperature combined with use of nickel in the fuel electrode

allows internal reform of hydrocarbon fuels within the fuel channels, obviating the

need for a separate fuel reformer. In addition, existing=fuel logistics systems can

be used with the MSOFC with essentially no modifications. Space-based systems that

are required to regenerate the hydrogen fuel from the reaction product water are

simplified by the fact that the MSOFC can use the same unit for power production and

fuel regeneration. The high operating temperature of the MSOFC allows the waste

heat to be rejected at a high temperature, thus minimizing the size of the heat

rejection system. The MSOFC systems are flexible in terms of their ability to

operate as an open cycle, closed cycle, or a combination of open and closed cycles.

Two types of systems are briefly described in the following paragraphs: (I) a

power system for terrestrial vehicle propulsion using hydrocarbon fuel, and (2) an

open-cycle, pulse-power system for use in space. Figure I is a schematic flow

dJa_ram of a 60 k_e vehicle propulsion power system designed to operate on a

methanol/water solution. The fuel is evaporated using residual heat from the spent:

fuel system. The incoming fuel flows through an eductor where a fraction of the

spent fuel is caused to recirculate to provide heating and sufficient water for the

reform reaction. The remaining spent fuel flows to the spent fuel burner where it

is combined with the air exhaust stream, and the heat from burning the spent fuel is

used to preheat the incoming air. The preheated air is further heated by

recirculating most of the hot air exhausted from the fue] ceil. This alr

recirculation is possible because very little Of the oxygen is Used in the: :

electrochemical reaction. The primary function of the air flow is to remove waste

heat from the fuel cell core. Reclrculating about two-thirds of the air minimizes

the size of the air preheater heat exchanger. In terms of size and weight of

components _ the fuel cell and air preheater are approximately equivalent, and the

other components are considerably smaller. All the components shown in this flow

diagram are enclosed in a high-temperature insulated box_ and other components such

as the fuel pump, air blower, and control systems are essentially at room

temperature,

The Volume power density of this system is approximately I kW/L, so the 60 kUe

system shown in figure 1 occupies about 60 L volume (-2 ft3). Operation of the

MSOFC system for vehicle propulsion will require tolerance to rapid changes in

thermal gradients. _hen a fuel cell operates at full powe}_ }he }em_erature of the

downstream region rises due to generation of waste heat. Operation Of the fuel cell

at low power produces a shallow temperature gradient because the quantity of waste

heat is low. Vehicle driving profiles typically require rapid shifting from high

power to low power and vice versa. Because the temperature gradients will change as

the power level changes, the MSOFC materials must be able to withstand rapid

temperature oscillations. The thin sections and honeycomb structure of the MSOFC

are particularly adapted to these rapid temperature changes. The coefficients of

thermal expansion of the four materials used in the MSOFC are well matched (ref. 5);

thus, changes in temperature do not cause significant stresses between the layers.

Coupling a nuclear reactor with a MSOFC appears attractive for space-based

pulse power systems. The MSOFC combines hydrogen and oxygen to produce the power

pulse, and the product water is stored on board. Following a test power pulse, the

42

m

2-

=

=

=

m

_m
fuel cell is run in reverse to regenerate the hydrogen and oxygen. A relatively

small nuclear reactor provides the power input for the reverse operation. Reject

heat from the reactor may be used to maintain the fuel cell in a "hot standby" mode

for rapid response to demands for power.

The hydrogen and oxygen are stored as gases in tanks at elevated temperatures

and pressures, as shown in figure 2. Emerging from the storage tank, the oxygen

stream passes through a blower and is preheated by mixing with an oxygen recycle

stream before entering the fuel cell. Similarly, the hydrogen passes through a

blower and preheater before entering the fuel cell. Within the fuel cell, DC power

is produced as oxygen is consumed and hydrogen is converted to water vapor. The

heat generated in the process is rejected as the oxygen recycle stream passes

through the radiator. After a reduction in temperature, the gaseous products are

stored. A sliding seal (or rolling diaphragm) is used to separate reactants and

products which are stored within the same tank.

The system can be operated in the reverse mode to regenerate the power

capability. The same monolithic fuel cell modules can be run backwards for the

regeneration: a separate electrolysis unit is not required. The length of time

required for regeneration is determined by the size of the external power source.

The overall efficiency of the system is attractive. The projected performance of

the system is to recover 72Z of the heat of reaction in the discharge mode and store

92% of the input electrical energy as chemical energy in the electrolysis mode. The

overall round-trip electrical efficiency is projected to be 66%.

The open-cycle concept is much simpler in system design and operation. A

simple once-through system that has excess hydrogen available for cooling the fuel

cell and for providing the chemical energy for the electrochemical process is shown

in figure 3. The waste heat is carried out of the fuel cell core with the hydrogen

stream, and the hydrogen and product water are dumped overboard. The oxygen stream

is recycled with only a small bleed stream to eliminate any impurities that might

accumulate. The system components include heat exchangers to heat the incoming

oxygen using waste heat from the fuel cell exhaust streams, an eductor to recycle

the hydrogen stream, and the fuel cell. Other components include the insulated

box, current collectors, and appropriate control systems.

This very simple system has application only in very special cases where very

high power densities are needed and an essentially "unlimited" source of hydrogen is

available. Specialized space power systems are the only currently known

applications for such a system, and this concept offers an interesting example of a

compact, high power density system for space. The power densities that can be

achieved with the MSOFC in this application are shown in figure 4. The fuel cell

power density is shown on the left y-axis, while the system power density is shown

on the right y-axis. The fuel cell power density is very high because the fuel

utilization is very low. Essentially pure hydrogen exists through the whole length

of the fuel channels, and the average Nernst potential is about i V. For these

calculations, the area-specific resistance was assumed to be 0.05 ohm-cm 2, which is

expected to be achieved with further development of the MSOFC. An important feature

of the system is the relationship between the fuel cell power density and the system

power density. As the fuel cell power density becomes larger, the contribution of

the balance-of-system begins to dominate the system weight. The fuel cell weight

becomes small relative to the balance-of-system when the fuel cell power density

exceeds about 20 kW/kg. For this system design, the system power density approaches

43
-4.5 kW/kg as an upper limit, and this power density is determined by the s_zes of

the heat exchangers required to heat the incoming hydrogen and oxygen and the

insulating box for the high-temperature components.

MSOFC FABRICATION

The principal challenge to fabrication of the MSOFC structure is to match the

sintering shrinkages of the four MSOFC materials: (i) the electrolyte, yttriastabilized zirconia, (2) the interconnect material, strontlum-doped lanthanum :

chromite, (3) the anode material, nickel-zirconia cermet, and (4) the cathode

material, strontium-doped lanthanum manganite. The MSOFC structure is assembled

while the tape-cast components are _n an unfired (green), flexible state. The ;_

polymer binder in the tapes bond the tapes together as they are placed in ContaCt "

and heated above the "glass transition" temperature of the polymer. The assembled

structure is then appropriately supported and heated slowly to burn out the polymer.

The structure is further heated to temperatures of about 1300 to _400°C to sinter

the materiais and create permanent bonds at the points of contact," It is important

that the shrinkages that occur during binder burnout and during sintering be

precisely matched among the four materials. This matching of shrinkages has been

accomplished satisfactorily during development of the MSOFC, and the results of this

shrinkage matching are reported elsewhere (ref. 5).

One of the important objectives in development of the MSOFC has been reduction

of the number of defects in the ceramic structure. Defects such as cracks allow the

fuel and oxidant gases to mix, thus reducing the efficiency and, if the crossleakage is severe, reducing the open-circuit potential. Defects such as non-bonds

between the layers increase§ the overall fuel cell resistance because of the reduced

cross-section of the conduction path. As the sintering shrinkages and the

coefficients of thermal expansion haye become better matched, the performance of the

fuel cells have improved. For example, the improvement in open-circuit voltage with

reduced number of cracks in the fuel cell is shown in figure 5. The large number of

cracks during the early development stages of the MSOFC caused sufficient crossleakage of fuel and oxidant to reduce the difference in oxygen potential across the

electrolyte. Since the Nernst potential is proportional to the log of the ratio of

oxygen partial pressure in the oxidant to that in the fuel, the open-circuit voltage

decreases with increased cross-leakage. The current status of MSOFC development is

such that essentially crack-free structures can be built and tested.

..... MSOFC PERFORMANCE

The performance of the MSOFC has app[ogched the theoretical limit as a result

of recent development efforts. The first MSOFC single cells had an area-specific

resistance (ASR) of about 8 ohm-cm 2, and Since that time the performance has

improved dramatically to the present state-of-the-art ASR of about 0.4 0hm-cm 2. The

theoretical limit of performance, based on materials resistances and elimination of

all electrode/electrolyte interfacial resistances, is about 0.093 ohm-cm 2 for an

electrolyte thickness of 50 _m. This ASR can be reduced to 0.05 ohm-cm 2 by reducing

the electrolytethickness to 25 _m. The interfacial resistances can be reduced with

further development (i) by decreasing the spacing between conductive particles in

the electrodes, (2) by increasing the thickness of the electrodes, and (3) possibly

by placing a mixed electronic/ionic conductor at the electrolyte/electrode

=

44

l
interfaces. These developments are expected to reduce the ASR to closely approach

the bulk resistance limit.

The performance of multi-cell MSOFC stacks is significantly less attractive

than the performance of single cells. Figure 6 shows the improvement in the

performance of multi-cell stacks since the operation of the first stack in

January 1986. While the performance improvement has been dramatic, it must be

improved further to equal that of the single cells. The improvement in open-circuit

potential is a result of improvements in fabrication methods, particularly in

matching the sintering shrinkages of the four different MSOFC materials.

The difference between single cell and multi-cell performance is primarily due

to interactions among the materials during the sintering process. A liquid-phase

sintering aid is used to assist densification of the interconnect material. This

sintering aid migrates throughout the MSOFC structure and affects the other

materials, particularly the cathode. The cathode becomes dense and tends to migrate

away from the electrolyte/cathode interface leaving voids in that critical area.

This effect on the cathode reduces the performance of the multi-cell stacks.

Future needs for development of the MSOFC include: (I) a more easily

sinterable interconnect material, (2) cathode materials that are more stable at the

sintering temperature, (3) development of ancillary materials such as insulation,

manifolds, cements, and current collectors, (4) optimization of MSOFC designs for

specific applications, (5) development of MSOFC systems, and (6) quality control

development, including improved control of source materials and development of

nondestructive analysis methods for the MSOFC structure.

CONCLUSIONS

The feasibility of the MSOFC concept has been proven, and the performance has

been dramatically improved. The differences in thermal expansion coefficients and

firing shrinkages among the fuel cell materials have been minimized, thus allowing

successful fabrication of the MSOFC with few defects. The MSOFC shows excellent

promise for development into a practical power source for many applications from

stationary power, to automobile propulsion, to space pulsed power.

ACKNOWLEDGMENT

This work was sponsored by the Defense Advanced Research Projects Agency of the

U.S. Department of Defense under ARPA Order No. 4850. Argonne National Laboratory

is operated for the U.S. Department of Energy under contract No. W-31-109-ENG-38.

i.

REFERENCES

Fee, D. C., Steunenberg, R. K., Claar, T. D., Poeppel, R. B., and

Ackerman, J. P., "Monolithic Fuel Cells," 1983 Fuel Cell Seminar

Abstracts, Courtesy Associates, Inc., Washington, DC, p. 74,

November 1983.

45
L

,

,

,

5,

McPheeters, C. C., Dees, D. W., Dorris, S. E., and Picciolo, J. J.,

"Argonne Monolithic Solid Oxide Fuel Cell Fabrication," 1988 Fuel Cell

Seminar Abstracts, Courtesy Associates, Inc., Washington, DC, p. 29,

October 1988.

Onoda, G. Y. Jr., and Hench, L. L., Ceramic Processing Before Firing,

John Wiley & Sons, New York, p. 411, 1978.

=

Richerson, D. W., Modern Ceramic Engineering, Marcel Dekker, Inc.,

New York, p. 206, 1982.

McPheeters C C. et al., ilRecent Advances in Monolithic Solid OxideFuel

Cell Development," 23rd Intersociety Energy Conversion Engineering ......

Conference, Vol. 2, Book No. I0272B, August 1988.

=

I The submitted manuscript has been authored I

I by e contractor of the U.S. Government I

under contract No. W-31-| 09-ENG -38.

I Accordin¢jly. the U. S. Government retains a

J nonexclusive, royalty-free license ,o publish |

or reproduce the published form of this

| contribution, or allow others to do so, for |

[ U. S. Government purposes.

= ....

±

46
0oLwo, 1

377 Solid Oxide I 365 13_

v_c.,,:o_v _ FuelCo, _ Ir_

/ 6e / II FUEL

AIR ll 274

377 scfm Air _ t"] Fuel

[] TB.,p.°C Preh°"te"l L-_I_;; 'I v"P°"z°rl

26.5 kwtl I Burn r I e.8 kWt I

o3m=i ,,,,_,,,, Olin=I

Air I_I tj 2g l

104 scfm

ORIGINAL PAGE IS

OF POOR QUALITY

lr_ 113_]Fuel, CH30H

Exhaust 23.4 kg/h

118 scfm 0.13 gpm

Figure i. Methanol fueled MSOFC for vehicle propulsion.

FUEL CELL SYSTEM RANKINE SYSTEM

.Y,Roo,.,.A,,.,......__ ,,o°"';"_7°_"" R...... I_

I

REACTOR PROOUCTE COOLER RAOIATOR

POWER CONDITIONER

SYSTEM ARRANOEMENT FOR CONTINUOUS. ALERT, AND BURST POWER

Figure 2. Space-based regenerative MSOFC system coupled

with a nuclear reactor for pulse power,

47
The MSOFC Power System Is:

1) Simple in Design

2) All Solid State

Figure 3.

TANK l ' J

! i

HYDROGEN_ MSOFC HYDROGEN

HIGH-TEMPERATURE ENCLOSURE

Open-Cycle System Schematic

Simple open-cycle HSOFC system with no hydrogen

regeneration and low fuel efficiency.

Monolithic Solid Oxide Fuel Cell

Open-Cycle System Power Densities

Can Exceed 4 kW/kg

45

40

•_ 35-

v

30-

"_ 25 -

20-

O 15-

-g

_ 10-

5-

10

9

13.

0

I I . I

0 1 2 3

Current Density, A/cm 2

Figure 4. Stack and system power densities for a simple

open-cycle space-based system.

48
EFFECT OF CRACKS ON OPEN-CIRCUIT VOLTAGE

1.2-

¢_ 1.1-

1- J

<

_ 0.9

_ 0.8

:3 0.7

0 .

0.6

a.

0 0.5

THEORETICAL OPEN CIRCUIT VOLTAGE

...... .o..... .... ... ...... ° .... o. .... .._8....= ..... . ... o... °°...°...°..oo. °--°-

o

0.4- I. I

0 10 2JO 30 40

CRACKS DETECTED BY X-RAY

Figure 5. Improved MSOFC open-circuit voltage with

reduction in ceramic defects.

IMPROVEMENT IN MSOFC STACK PERFORMANCE

1.2-

> 1.0.

0.8

LtJ

_ O.6

Q.)

_ 0;4

<

,_" 0.2-

MAXIMUM CELL PERFORMANCE

-'° ........................................................... o.....°°o°.o.°H.

- _ FEBRUARY 198_ o

0.0-

0 10 20 30 4'0 50

CURRENT DENSITY, mA/cm 2

Figure 6. Improvement in multi-cell HSOFC stacks during

development at Argonne.

49
i|!||

mE i!=m

,,ZLzz
N90-20461

THE ELECTROCHEMICAL GENERATION OF USEFUL CHEMICAL

SPECIES FROM LUNAR MATERIALS *

Kan J. Tsai, Daniel J. Kuchynka and Anthony F. Sammel_s

Eltron Research, Inc.

Aurora, Illinois 60504

Electrochemical cells have been fabricated for the simultaneous generation of

oxygen and lithium from a Li20 containing molten salt (Li20-LiCI-LiF). The cell

utilizes an oxygen vacancy conducting solid electrolyte, yttria-stabilized zirconia

(YSZ), to effect separation between oxygen evolving and lithium reduction half-cell

reactions. The cell, which operates at 700-850°C, possesses rapid electrode kinetics at the lithium-alloy electrode with exchange current density (i o) values being

>60mA/cm 2. When used in the electrolytic mode, lithium produced at the negative

electrode would be continuously removed from the cell for later use (under lunar conditions) as an easily storable reducing agent (compared to H 2) for the chemical refining of lunar ores. Because of the high reversibility of this electrochemical

system, it has also formed the basis for the lithium oxygen secondary battery system

which possesses the highest theoretical energy density yet investigated.

INTRODUCTION

The strategy being pursued for lunar ore refining is based upon electrochemical

cells possessing the general configuration:

Li alloy/Li20-LiCi-LiF molten salt/YSZ/Lao.89Sro.11MnO 3

In practical electrolytic cells, lithium produced at the negative electrode would,

after removal from the cell, be available for lunar ore refining via the general

chemical reaction:

2Li + MO _ Li20 + M

where MO represents a lunar ore. Emphasis to this time has been on Fe203, TiO 2 and

the lunar ore ilmenite (FeTiO3), all of which have been shown chemically reducible by

Li to give metals. The resulting Li20 reaction product could then be removed from

the solid-state reaction mixture by sublimation and reintroduced into the negative

electrode compartment of the electrolytic cell to be electrolyzed again. Hence, this

electrochemical approach provides a convenient route for the simultaneous generation

of both metals and oxygen from lunar materials on the Moon's surface. It has previously been suggested that oxygen might be extracted from ilmenite by its initial

chemical reduction by hydrogen, initially transported from Earth (ref. i-3), and presumably requiring cryogenic storage on the Moon's surface. Other approaches discussed for the chemical reduction of lunar ores have included carbothermic reduction

to give the desired metal (ref. 4,5). Previous work by others has also investigated

the direct high temperature electrochemistry of simulated lunar materials (molten

*This work was supported by NASA Johnson Space Center under Contract No.

NAS9-17991.

5]

PRECEDING PAGE BLANK NOT FILMED
silicates), using platinum electrodes. Electrolysis of these melts demonstrated simultaneous oxygen evolution at the anode (ref. 6-10) and deposition of an impure

metal-silicon alloy slag at the cathode (ref. 11-14). Nevertheless, this preliminary

work showed feasibility for high temperature molten salt electrochemistry as an attractive approach for oxygen evolution and metal reduction from lunar type ores. In

the approach being discussed here, oxygen evolution and lunar ore reduction are separated into distinct electrochemical and chemical steps. This strategy minimizes the

opportunity for lunar ore originating impurities from entering the electrochemical

cell and degrading performance.

The reducing power of lithium is significantly greater than that of hydrogen.

This fact can be summarized by the two general reactions:

H 2 + _O 2 _ H20 AG = -68kcal/mole (25°C) (1)

and

2Li + _02 _ Li20 AG =-142kcal/mole (25°C) (2)

The corresponding free energy of reaction between lithium and candidate metal oxides

of interest at 1000°K are summarized in Table i.

TABLE 1

4Li + SiO 2 + 2Li20 + Si

2Li + FeO + Li20 + Fe

4Li + TiO 2 + 2Li20 + Ti

6Li + Cr203 + 3Li20 + 2Cr

AG(kcal/mole)

-28.0

-53.9

-15.6

-87.9

We estimate the free energy change associated with ilmenite reduction by lithium at

IO00°K, corresponding to the overall reaction:

FeTiO 3 + 6Li _ Fe + Ti + 3Li20 (3)

to be AG = -62.2kcal/mole. Hence, lithium would appear to be a sufficiently strong

reducing agent for many of the transition metal oxides anticipated to be present on

the Moon's surface.

The electrochemistry occurring in these cells possesses high reversibility and

consequently has given us the opportunity to study the lithium/oxygen secondary battery. The overall electrochemistry occurring in this cell upon discharge/charge

cycling can be represented by:

discharge

FeSi2Li x + _02(air ) = m ei20(in molten salt) + FeSi 2 (4)

charge

with overall capacity being dictated by the concentration of Li20 that may be incorporated within the negative electrode compartment of this electrochemical cell. Although theoretical energy densities for a given battery system are often a poor indicator of the final practical energy density that might be achieved, it is of interest

to note that the energy density for electroactive materials in the lithium-oxygen

system calculates to 4082Wh/kg compared to 2211Wh/kg for the lithium-chlorine cell;

the highest energy density secondary system investigated to this time. The significant technical observations to be discussed which make this an attractive secondary

battery system are i) both electrodes possess high electrochemical reversibility,:

52

- 7

: 7 <
li) the anodic reaction upon cell charge involves the exclusive evolution of oxygen,

and iii) a significant concentration of Li20 can be incorporated into the Li20-LiCILiF ternary melt. A comparison of the lithlum-oxygen secondary cell with other major

conventions and advanced systems is shown in Table 2.

TABLE 2

Major secondary batteries for electrical energy

storage in comparison to the lithium oxygen system

Battery System Negative Positive OCP(V)

Lead-Acid Pb PbO 2 2.0

Nickel-Cadmium Cd NiOOH 1.2

Nickel-Zinc Zn NiOOH 1.7

Zinc-Chlorine Zn CI 2 2.12

Lithium-Iron

Sulfide Li(AI) FeS 1.2

Sodium-Sulfur Na S 2.1

Lithium-Chlorlne Li C12 3.53

Lithium-Oxygen Li 02 2.3

Theoretical Energy Density

for Electroactive Matls. (Wh/k 8)

238

148

448

829

626

760

2211

4082

RESULTS AND DISCUSSION

Previous work performed in our laboratory (ref. 15) showed that the ternary molten salt system Li20-LiF-LiC1 is an attractive candidate electrolyte for this electrochemical cell. The utility of this molten salt electrolyte was dependent, however,

upon gaining some insight into the amount of Li20 that could be accommodated in the

LiF-LiCI (70m/o) molten salt and its impact upon the resulting solid-liquidus curve.

This was performed by progressively introducing Li20 in 2-5m/o increments into a LiFLiCI (70m/o) binary mixture in a conventional conductivity type cell. After each

Li20 addition, the ternary molten salt mixture was initially heated to 550°C, at

which temperature it became molten. The resulting molten salt ionic conductivity

values were measured between stainless steel electrodes (icm 2) placed 1.5cm apart

during slow cell cooling so as to minimize interelectrode thermal gradients. The

molten salt freezing point was detected by a dramatic decrease in measured ionic conductivity. The resulting solid-liquidus curve obtained for this ternary system is

shown in Figure i, which indicates that for Li20 contents above ml8m/o the molten

salt goes into a two-phase region consisting of Li20 (18m/o),LiF (24.6m/o),LiCl (57.4

m/o) in direct contact with solid Li20. The important point here is that excess Li20

could be introduced into the electrolyte cell and would not promote melt freezing.

Both half and full cell electrochemical measurements have been performed on this

cell. Coulometric lithium deposition has been performed on both iron and iron silicide (FeSi2) negative electrode substrates. Iron silicide was of interest since

lithium deposition results in formation of a series of distinct electrochemically reversible lithium compounds up to the composition FeSi2Lil0. This has not only permitted us to conveniently store lithium in the negative electrode compartment for the

reversible lithium oxygen storage battery, but also acts as a substrate for wetting

lithium, facilitating its later removal from the negative electrode compartment of

the electrolytic cell. Preliminary current-potential characteristics of this electrode were performed in electrolytic cells possessing the general configuration shown

in Figure 2. The charging current-potential curves for this cell (Figure 3) clearly

show voltage plateaus associated with each ternary lithium compound. Electrode ki5S
Figure i.

500

3

&SO

X 4_0

441)

o

Y 420

4O0

f liquid liquid

r-l,i20-l-iF-1,1C11 LI20-LIF-LIC] + Li20 solid

I

"-_L _ ......

solid

Li20-LIF-LICI

I0 20 30 40 50 60

ORIGINAL PAGE IS

OF POOR QUALITY

Li20 Mole Percentage Introduced into LIF-LiCI (70m/o)

Solid-liquidus curve for Li20-LiF'LiCI molten salt as a

function of Li20.

i

Af

...... : 7 2c i

f_\iii : i:::

Figure _. _Sdhematic drawing of cell configuration used for the Simultaneous electrolytic generation of lithium and oxygen.

A) anode current collector; B) caicia or yttria stabilized

zirconia; C) molten salt; D) Lao.89SrO.lIMnO 3 anode; .....

E) stainless steel or FeSi 2 cathode; F) furnace.

netics for this electrode were performed using a molten salt of composition Li20

(6.4m/o),LiF(28.1m/o),LiCl(65.5m/o) using FeSi2Li x for the working and SiLl 4 for

counter and reference electrodes. For voitage plateaus corresponding to the respective ternary alloys FeSi2Li6, FeSi2Li 8 and FeSi2Lilo cyclic voltammetry scans

(50mV/s) were performed ±200mV from its initial open circuit potential (OCP) with

respect to the SiLi 4 reference. A representative current-overpotential curve is

shown in Figure 4, which was analyzed by use of the Allen-Hickling relationship:

54
ORIGINAL PAGE iS

OF POOR QUALITY

t _nF (5)

log _nF/RT = log io 2.303RT

e - 1

where t o was the exchange current density from which a direct estimation of relative

electrode kinetics were determined, _ was the electrode overpotentlal, _ the transfer

coefficient and the other symbols had their usual significance (ref. 16). Exchange

current densities (i o) were obtained from the intercept by plotting _ versus

i

log _nF/RT The dependency of io upon both cell temperature and lithium

e - i

alloy composition is summarized in Table 3. These current-voltage measurements

clearly showed the presence of rapid electrode kinetics at the negative electrode

suggesting that this Faradaic process will not be the rate limiting process in the

finally developed electrochemical cell.

Figure 3.

2.3

2.2

-_ 2.1

2.o

._ 1.9

[.7

FeSI2LI 6

FeSi2Li4/_

FeSi2],1 I0

I

lO0 150 200 250 300

Time (mln)

IR free charge curve for FeSi 2 in 28.5m/o LIF, 66.5m/o LiCI

and 5m/o Li20 at 650°C vs La0.89Sr 0 llMnO3/Pt (air). Current

density 10mA/cm 2 .

Cathodic Current Density

mA/cm 2

'i

Anodie _D_ ,to loo _o Cathodic

Overpotentlal ÷ ---_---4---_ -_--_--_----+--+ Overpotentlal -_0 -_oo -i_0 -r00

_,(_v) il n(mV)

Anodic Current Density

mAlcm 2

Figure 4. Current-overpotentlal curve for the cell

FeSI2LII0/Li20-LIF-LICI/SiLI4

at 800°C. Initial open clrcult potential -71mV with respect

to SILl 4 reference. Scan rate 50mV/sec.

55
TABLE 3

Cell Temperature Measured Exchange Current Density io (mA/cm 2)

°C FeSi2Lilo FeSi2Li 8 FeSi2Li 6

650 49.7 50.6 43.7

700 54.5 54.7 50.9

750 56.0 57.2 55.6

800 64.5 65.7 58.2

In fully developed electrolytic cells we must facilitate the convenient removal

of faradaically deposited lithium at the negative electrode. One of the major obstacles to achieve this is related-to the high surface tension of Li at 700°C (320

dynes/cm) which compares to Hg (400 dynes/cm) at room temperature. Deposited lithium

resident in a porous electrode structure, in the absence of wetting, dan as a conse _

quence prove difficult to phYSically remove via a proximate orifice. We have found,

however, that both SiLi x and FeSi2Li X can act as effective substrate sites for wetting faradaically deposited lithium, thereby providing a strategy for the subsequent

continuous removal of lithium from the electrolytic cell. A schematic of the experimental arrangement used for c0mparing approaches for lithium removal is shown in Figure 5. Lithium could be conveniently removed from the cell by application of mbpsi

into the lithium chamber C, thereby enabling its transfer into glass trap D. In the

event of lithium/molten salt mixtures being formed in the electrochemical cell, we

have found that effective segregation into two distinct phases can be achieved if the

stainless Steel transfer tube between the electrolytic cell and lithium storage compartment is initially silicided to form a thin coating of FeSi 2. Subsequent reaction

of this thin coating with lithium to give FeSi2Li x will provide the necessary wetting

characteristic necessary for promoting preferential lithium migration from the electrolytic cell.

ORIGINAL PAGE iS

OF POOR QUALITY

I/8" tube

lfl l I

jc

porous SILl x

alloy

_ vacuum

Figure 5. Schematic diagram for simulated lithium removal process.

C) lithium containing chamber; D) lithium collecting chamber;

F) furnace; L) lithium; P) pressure gauge

As we have previously discussed, lithium is a strong reducing agent which should

in principle, be able to reduce simulated lunar ores to the corresponding metal.

This has been examined by investigating direct reaction between lithium and the respective metal oxides TiO 2, Fe203 and FeTiO 3. Here, thermal analyses were performed

56

=

=-_

L

w
under both argon and vacuum (~i00_). Lithium and the metal oxide of interest were

initially placed into a small boron nitride crucible located inside a 3/4" Swagelock

union used as the reaction chamber. A chromel alumel thermocouple was placed in direct contact with the reaction mixture. Because of the anticipated high reducing

ability of lithium, only 0.02g was used for each reaction with an excess of metal

oxide being present. The reaction chamber was heated at 5°C/min. The inception of

solid-state reaction was manifested by a rapid temperature increase in all cases investigated, indicating the reaction to be highly exothermic and irreversible. Thermal analyses for respectively Fe203 and TiO 2 are shown in Figure 6 A and B. The reaction product between Li and Fe20 3 was clearly shown to be elemental iron and could

be conveniently removed with a magnet. Analysis of the reaction mixture between Li

and TiO 2 was determined by initially dissolving with 37W/o HCI. Here, unreacted TiO 2

was insoluble, where Ti became dissolved as TiC13 to give a dark blue solution. Quantitative determination of elemental Ti was then performed spectrophotometrically.

Based upon the initial Li present in the reaction mixture, approximately 80% participated in promoting TiO 2 reduction to Ti. Temperatures experimentally observed for

the inception of these reactions are summarized in Table 4 together with the corresponding calculated heat of reaction at each respective temperature.

ORIGINAL PAGE iS

OF POOR QUALIFY

Figure 6.

994

509

391

o 271

$

147

,w

25

745

626

509

391

271

147

25

_her_sl excursLon

corrempo.dt.g [o --_

reactton

6Lt + V. 203 * 3L120 ÷ 2re

therma I excursion

_or re spond1rig co

_+ • . ,

0.5 1 1.5

Time, Hr

Thermal analysis for detection of solid-state reaction

temperature between Li and Fe20 3 (excess) A, and TiO 2

(excess) B, under argon.

This electrochemical technology, as we have already discussed, is also compatible for a high energy secondary battery. Initial work (ref. 17) towards this goal

focused upon electrochemical cell configurations analogous to that shown in Figure 2

57
ORIGINAL PAGE iS

OF POOR QUALITY

TABLE 4

Experimentally observed temperatures and heat of reaction for the inception

of reaction between lithium and the metal oxide of interest.

Temperature for Inception Heat of Reaction kcal/mole,

Metal Oxide of Solid-State Reaction, °C Calc'd at Temperature

TiO 2 457 -71.6

Fe203 440 -244.8

FeTiO 3 (ilmenite) 432 --

Fe203+TiO 2 (i: !) 428 -157.2

where maintenance of an inert (Ar) atmosphere in the negative electrode component

:Could be conveniently achieved (ref. 18). A preliminary Charge'disCharge curve fo E

this type cell using a FeSi2Li x negative electrode is shown in Figure 7. Again, the

voltage plateaus Observed correspond to the reversibi$_ormation of ferro silicon

lithium ail0ys. Furthermore, it was found during cell charge that the volume of oxygen generated was in fact Faradaic, demonstrating that even in direct contact with

halide containing molten salt, the solid electrolyte remains an exclusive oxygen

anion conductor.

2._0

2, O0

CHARGE DI 5CF, ARG E

Figure 7.

0 5 10 !5 ?O

Time (Hours)

IR-free charge-discharge curve for the cell LixFeSi2/(52.5-

23.6-23.9)m/o LiCl-LiF-Li20/ZrO 2 (5W/o CaO)/Lao. 89Sr0.11MnO3 /

Pt at 20mA/cm 2 (at negative electrode). Total cell resistance

24_. Temperature 650°C.

For practical reversible lithium/oxygen batteries for terrestrial applications,

however, we must identify a design which will enable us to operate in the atmosphere.

This has been addressed by fabrication of small prototype cells possessing the gen- z

erai design shown in Figure 8. The design is based upon a 5cm high YSZ crucible possessing respective inside and outside diameters of 2.1 and 2.35cm (totai voiume 19ml). -Z

Cells were fabricated_Inthe partially charged state using FeSi2Li 8 as the negative

electroactive material. Current collection from this electrode was via an aluminized

nickel chromium aiioy_(Ni (76%), Cr (16%), AI (4.5%), Fe (3%), Y (trace), Hayn@s --"

Alloy 214). The La0.89Sr0.11Mn03 oxygen electrode was initially prepared (ref. 19)

from a 15W/o suspension in polyvlnyl alcohol-ethylene glycol of the appropriate stoichiometric metal nitrates. Current collection from this region was via a Pt wire

initially tightly coiled in this area. Synthesis of the perovskite electrocatalyst

was achieved by heating in the atmosphere to _I000°C for lh. Protection of the negative electrode compartment from direct contact to the atmosphere was achieved using a

machined alumina cover (Cotronics, Inc.):whose residual porosity was removed by'use

of a coating of high density ceramic cement (Sauereisen #8), followed by curing at

58
Figure 8.

C

ORIGINAL PAGE iS

OF POOR ^"^_ _'rv

Schematic design of the lithium/oxygen secondary storage

cell discussed in the text. A) current collector nickelchromium (80:20) (-); B) platinum current collector (+)

tightly wound around YSZ crucible; C) machined alumina cap;

D) YSZ 02- conducting solid electrolyte; E) zirconia felt

separator; F) molten salt containing 66.5m/o LiCI, 28.5m/o

LiF and 5m/o Li20; G) Lao.89Sro.11MnO3 electrocatalyst;

H) FeSi2Li 8 negative electrode

90°C for one day.

In partially charged cells fabricated here, the negative electrode compartment

initially contained a molten salt of composition LiCI (66.sm/o),LiF (28.sm/o),Li20

5m/o). This cell was operated in the atmosphere by initially heating to 600°C, at

which temperature the open-circuit potential became 2.2V. Between 650 ° and 800°C

the overall cell resistance decreased from 47_ to 3.9_. This reflected the improved

02- conductivity for the solid electrolyte at higher temperatures. Ea was found to

be 16kcal/mole, similar to literature values (ref. 20). Current-potential and discharge curves for this cell are shown in Figure 9 (A and B).

2.5

::>

-._ 2.0

,--4

a_ 1.5

.,-t

1.0

.u 0.5

o

o

A

2.70

::>

2.20

0_ 1.70

4J

O

::>

[ .20

i--(

c._ 0.70

B DISCHARCE

l t I i J l

0 5 10 15 20 o 1 2 3

Current Density (mA/cm 2) Time (hr)

Figure 9. Current-potential and discharge curve for the lithium/oxygen

cell at A) 800°C (Rcell - 3.9_) and B) 850°C at a C/50 rate

(5.3mA/cm2).

The above discussed results demonstrate that this electrochemical technology is

evolving into a practical option for both the generation of chemical species on the

59
Moon's surface and as a high energy secondary battery for either terrestrial or lunar

applications.

REFERENCES

i. H. P. Davis, Eagle Engin. Report #EEl 83-63, Houston (1983).

2. M. A. Gibson and C. W. Knudsen in "Lunar Bases and Space Activities of the 21st

Century," NASA, 1984, p. 26.

3. E. Kibler, L. W. Taylor and R. J. Williams in "Lunar Bases and Space Activities

of the 21st Century," NASA, 1984, p. 25.

4. W. C. Phinney, et. al., in "Space Based Manufacturing from Nonterrestrial Materials, Amer. Inst. of Aeronautics and Astronautics," New York (1977).

5. A. H. Cutler in "Lunar Bases and Space Activities of the 21st Century," NASA,

1984, p. 22.

6. K.W. Semkow and L. A. Haskin, Geochim. Cosmochim. Acta, 49, 1897 (1985).

7. K. W. Semkow and L. A. Haskin, Abstracts to Lunar Planet Sci., XVl, (1985)

p. 761.

8. A. Ghosh and T. B. King, Trans. Metall. Soc., AIME, 245, 145 (1969).

9. J. K. Higgins, Glass Technol., 23, 90 (1982).

i0. J. K. Higgins, Glass Technol., 23, 180 (1982).

11. D. S. Kesterke, U.S. Dept. of Interior Bureau of Mines Report of Invest. RI-7587,

1971.

12. D. J. Lindstrom and L. A. Haskin in "Space Manufacturing Facilities 3," T. Grey

and C. Krop, Eds., Amer. Inst. Aeronautics and Astronautics, New York, 1979.

13. W. F. Carroll, Jet Propulsion Laboratory Report 83-36, 1983.

14. R. H. Lewis, D. J. Lindstrom and L. A. Haskin, Abstracts to Lunar Planet Sci.,

XVI, 1985, p. 489 .......

15. A. F. Sammeils and K. W. Semkow, J. Power Sources, 22, 285 (1988).

16. P. A. Allen and A. Hickling, Trans. Faraday Soc., 53, 1626 (1957).

17. K. W. Semkow and A. F. Sammells, J. Electrochem. Soc., 134, 2084 (1987).

18. K. W. Semkow andA._ F. Sammells, J. Electrochem. Soc., 134, 2088 (i987).

19. H.-M. Zhang, Y. Teraoka and N. Yamazoe, N. Chem. Lett., 665 (1987).

20. T. H. Etseii and S. N. Flengas, Chem. Rev., 70, 339 (1970).

m

z

z

i

60

=

R
N90-20462

A NEW CONCEPT FOR HIGH-CYCLE-LIFE LEO:

RECHARGEABLE MnO2-HYDROGEN

A.J. Appleby, H.P. Dhar, Y.J. Kim and O.J. Murphy

Center for Electrochemical Systems and Hydrogen Research

TEES-Texas A&M University

College Station, Texas 77843

The nickel-hydrogen secondary battery system, developed in the early 1970s, has become the

system of choice for geostationary earth orbit (GEO) application. However, for low earth orbit (LEO)

satellites with long expected lifetimes the nickel positive limits performance. This requires derating of

the cell to achieve very long cycle life. A new system, rechargeable MnO2-Hydrogen, which does not

require derating, is described here. For LEO applications, it promises to have longer cycle life, higher

rate capability, a higher effective energy density, and much lower self-discharge behavior than those of

the nickel-hydrogen system.

INTRODUCTION

Space power systems designed for continuous high-power-output applications have most

commonly been based on the use of photovoltaic panels. Since in all orbits satellites are subjected to a

greater or lesser number of eclipses, a reliable means of energy storage and power load-leveling is

required using reliable, maintenance-free secondary batteries. In the early 1970s, nickel-cadmium was

the only proven long-lived secondary battery for use in space. Most of its long-term applications were

in GEO communications satellites, for which there are two 45-day eclipse periods per year and one

eclipse per eclipse-day in typical orbits. The eclipses are of variable duration, and reach a maximum of

1.2 h at the middle of the eclipse period. Over a ten-year satellite lifetime, a total of about 900 cycles

are therefore required, varying from a maximum of medium discharge (60% depth of discharge, DOD)

at the C/2 rate for 1.2 h, to shallow discharge, 60% DOD being a precautionary choice for long lifetime

and reliable operation. Typical aerospace Ni-Cd ceils can easily exceed 2000 cycles under these

cycling conditions, but in a battery their energy density is little better than 20 Wh-kg -1. In view of the

importance of battery weight in determining overall power system weight, hence incremental satellite

structural weight, new battery concepts offering higher specific energy were explored. As an

illustration of the battery contribution to power system weight, state-of-the-art deployable photovoltaic

(PV) panels have a specific power of 55 W-kg -1, whereas the Ni-Cd battery under the above discharge

conditions has a specific power of 16.7 W-kg -1. Thus, 18.2 kg of PV array is required to produce a

continuous 1 kW output, with an additional 1.2 kg (three-axis stabilization) to charge an 80% efficient

battery over 22.8 h per eclipse-day, discharging in 1.2 h with a power output of 1 kW. Based on

16.7 W-kg -1, such a battery would weigh 60 kg, thus demonstrating the driving force for reducing

battery weight.

HIGH-ENERGY-DENSITY SATELLITE BATTERIES

The first attempt to reduce battery weight was the development of a regenerative fuel

cell/electrolyzer (RFC) in the 1960s [1]. This consisted of a thin, lightweight pressure vessel

containing pure noble metal negative and positive electrodes of fuel cell type, with a separator

containing KOH electrolyte. A flexible sealed divider cut the pressure vessel into two volumes in the

6]
ratio 2 • 1 to storehydrogenproducedat the negativeandoxygenat thepositive. The gaseswere

recombinedat thesameelectrodes.The systemshouldbecapableof manydeep-dischargecycles,

with indefinitecycle-life atthehydrogenelectrode,andlonglife attheoxygenelectrodeprovidedthat

suitableelectrocatalysts(e.g.,gold-iridium)areused.A calculationshowsthat54Wh/kg isobtainable

[2], assuminganInconel718pressurevesselwith a safetyfactorof 4. Thisis overtwicethevaluefor

Ni-Cd. However,thestate-of-the-artelectrocatalysttechnologycouldnotoffer around-tripefficiency

greaterthanabout50%dueto oxygenelectrodepolarization. More importantly,thestoragesystem

was subject to cross-leaks,which limited reliability, lifetime, managementof the excesswater

producedon dischargeand consumedon charge,andproved to be a problem. In spite of these

drawbacks,anupdatedversionof this systemusingmorerecenttechnologymaystill haveadvantages

for certainapplications[2]. _

To avoid cross-leaksandelectrolytemanagementproblems,attentionwastransferredto metaloxygenbatteryconcepts. Thesecanin principle show high energydensity(for example,the zincoxygen couple has a practical energy density of 600 Wh-kg-l), and the open volume of the pressure

vessel (about 60% of the weight of the RFC) is small, since only oxygen storage is required. In

consequence, zinc-oxygen systems (which do not have a water-mangement problem since the dicharge

product is simply zinc oxide) were estimated to have practical energy densities on the order of 90 Whkg -1 [3]. However, like all alkaline zinc systems in stationary electrolyte, ihe negative showed

dendrite formation and shape-change, resulting in limited cycle-life. The zinc electrodecan aisoreact

explosively with excess oxygen on overcharge, which is accompanied by drying out. Finally, while

the coulombic efficiency of the zinc-oxygen system is good, the oxygen electrode charge-discharge

characteristics limit voltage efficiency to 60%. Since the corresponding cadmium system [3] has only

the advantage of cycle life over zinc-oxygen, with a considerably poorer energy density, attention

focused on metal-oxide/hydrogen systems, specifically nickel-hydrogen, which combined the best

electrodes of the Ni-Cd and RFC.

THE NICKEL-HYDROGEN SYSTEM

The Ni-H2 battery combined the two electrodes from the Ni-Cd battery and the RFC with the best

electrochemical characteristics. Both operate close to reversibility, and give high coulombic

efficiency. The battery was first described in the patent literature [4], and was developed from 1970

onwards, primarily at Comsat Laboratories [5]. No attempt was made to use specially-developed

components for the battery. Compared with zinc-oxygen, it was immediately obvious that the nickel'

hydrogen system would be conservative from the viewpoint of energy density. Aerospace Ni-H2

positives of conventional construction have a typical capacity of 0.11 Ah/g, which is perhaps only

20% of that of a zinc negative, and hydrogen negative systems require twice the container volume and

weight than those using oxygen positives. While a lightweight design, with a minimum container

safety factor, might achieve 55 Wh-kg -1, a more practical figure is 45 Wh-kg -1 on the Single-cell level

[5]. _: _ _ _ _

The Ni-l_2 Systemhas no problems of electrolyte management; since the overall ceU reacti0n

[NiOOIY¥ 1]2H2 _ Ni(OH)2] does not involve formation or consumption of free water. This

represents an advantage over both the RFC and the Ni-Cd system. The hydrogen electrode operates

under benign conditions, at high hydrogen pressures (up to several hundred psi) and with maximum

current densities of about 40 mA/cm 2 at the C rate. There is no evidence of change of hydrogen

electrode performance with time under normal operating conditions. While the nickel electrode has an

enviable record of longevity, it is still the life-limiting element in the battery. Since the nickel electrode

operates in a starved electrolyte environment, any electrolyte loss from the separator can cause cell

failure. With chemically-impregnated aerospace electrodes, failure occurred at about 2500 cycles

under GEO cycling conditions with 50% overcharge per cycle. Corrosion of the nickel sinter, with

consequent consumption of water, was identified as the cause of failure [6]. Improved lifetimes (to

about 5000 cycles) were observed with under-impregnated electrodes, and with electrochemically

62

i

!

=

=

E

z

=

E

r.

=

=

E

c_
impregnated electrodes [7], which have a more uniform distribution of active nickel hydroxide in the

sinter structure [6].

On cycling, swelling of the nickel electrodes occurs, resulting from volume changes in the active

material. Corrosion of the sinter in effect produces more active material, but the internal stresses

created cause progressive breakdown of the sinter by the rupture of the "necks" between individual

nickel particles. Since the electronic pathways through the sinter are thereby compromised, capacity

slowly falls after going through a maximum early in life. Eventually, breakdown is total, and the cell

ceases to function, even if electrolyte losses are made up. The situation is improved by using lower

active material loadings, and by operating the electrodes at lower depths of discharge, keeping

overcharge to a minimum. This demonstrates the effect of rate on cycle life.

For LEO applications, charge/discharge rates are not only much higher than those for GEO, but a

far larger number of cycles is required. In an equatorial orbit, sunlight is 0.9 h, and the eclipse period

is 0.6 h, with one charge-discharge cycle per orbit, or 5840 cycles per year. Consequently, for a

guarantee of a ten-year lifetime, over 100,000 cycles would have to be reliably demonstated in

laboratory testing. Thus far, 8000-12,000 cycles can be routinely demonstrated with low-loading,

electrochemically impregnated electrodes. One aspect of recent work has been the effect of initial

electrolyte concentration on cycle-life. It has been shown that life is at a maximum at a starting

concentration of 26% KOH, when about 30,000 LEO cycles can be achieved with derated electrodes

[8]. This seems to suggest that two different processes control electrode lifetime, one perhaps being

charge acceptance of the nickel positive, which is favored at higher electrolyte concentrations, and the

limiting effects of sinter corrosion, which also become more marked at high concentrations. However,

in spite of this work, it would appear that the nickel electrode, at least in its present form, is marginal

for long-term LEO use.

THE RECHARGEABLE MnO2 POSITIVE

Electrolytic manganese dioxide has only a limited charging ability (about 50% of the first electron)

[9]. However, in recent years a highly rechargeable form of MnO2 has been demonstrated, in which

both electrons can be reversibly charged and discharged. This consists of a material doped with

bismuth (1-5 atom %), with birnessite structure [10-12]. It can be subjected to many thousands of

rapid cycles (at the 12C rate in thin electrodes) and still shows no noticeable degradation in perfomance

(Figure 1). It is quite indifferent to overcharge (no formation of soluble species) and to overdischarge.

Unlike electrolytic MnO2, both electrons discharge in the same potential range, giving a very flat

discharge plateau and a more useful capacity to the voltage cut-off point, as Figure 2 shows for a cell

with a zinc negative. In a benign environment (i.e., in the absence of certain impurities in the

electrolyte) 85% of the two-electron capacity can be recovered over an indefinite number of cycles.

Like Ni-H2, MnO2-H2 has no water mangement problems, since the overall cell reaction is MnO2 +

H2 _ Mn(OH2). Finally, overall efficiency is estimated to be 78% (similar to Ni-H2), and heat

dissipation has similar characteristics to those of Ni-H2.

At first sight, the proposal of using the modified MnO2 positive combined with a hydrogen

electrode does not make much sense, since the couple has a comparatively low voltage compared with

that of nickel hydrogen (average plateau potential at LEO rates is approximately 0.7 V versus hydrogen

at an average pressure of 24 atm). However, the low potential has certain operational advantages that

are conducive to long life and low corrosion rates. In addition, since the modified MnO2 positives can

be operated at a nominal 85% of their two-electron capacity in long-term cycling, even at high rates,

the positives are much lighter and have smaller volumes than nickel electrodes. Finally,

microcalorimeter measurements (Figure 3) have established that the rate of the self-discharge reaction

between MnO2 and compressed hydrogen is much less than that for NiOOH. The mild corrosion

environment allows the use of graphite powder additive as an electronic conductor in Teflon-bonded

electrodes. With a material doped with 2 atom % bismuth, containing 15 wt % graphite and 7 wt %

Teflon binder, the specific coulombic yield of the positive mix is 0.43 Ah/g to 85% DOD. Applied to a

screen similar to those used for hydrogen electrodes, this value is reduced to 0.249 Ah/g.

63
COMPARISONBETWEENMODIFIED MnO2-H2AND Ni-H2FORLEO USE

Table 1 showsa breakdownof componentweightsin a Ni-H2 cell with a nominal capacityof

38.5Ah (to 1.0V at23°C). Thevolumeof theInconel718container(with safetyfactorof 4) hasbeen

arrangedtOC6nt_n38.5Ah of hydr0genbetween41and7 atm. To ensurevery long=cyclelife Under

LEO conditions,thepositiveelectrodesin suchasystemmustbederated,typicallyto 40%of capacity.

Thisrequiresa largerstackweightandvolume,sincethenumberof positives,negatives,separators,

andgasstreams,togetherwith thetotalelectrolytevolume,mustbeincreasedby 2.5. In consequence,

thecontainerweight andvolumemustbecorresopondinglyincreased.Overall,theLEO Ni-H2 cell

hasaneffectivespecificenergyof 22.2_g (33.I Wh/1),whereasthebaselineNi-H2 hasvaluesof

44.9Wh/kg and51Wh/1,

The final coIumncontainsthecorrespondingbreakdownfor LEO Mno2-H2. It can be seen that

the smaller weight and volume of the positives results in a lighter and more compact cell than baseline

Ni-H2, although its specific energy is degraded by its lower operating voltage. However, in the LEO

application it offers a specific energy of 32.6 Wh/kg, almost 50% higher than the purpose-designed

Ni-H2 cell with derated electrodes, intended for the same application. On a Wh/1 basis, both systems

have the same practical energy density.

CONCLUSIONS =

The proposed modified MnO2-H2 cell operating at 0.7 V (nominal), with electrodes operating at

85°2Uoftwodectron capacity has the Same volume energy density as Ni-H2 operating with positives

derated to 40% of capacity to achieve 25,000 or more LEO cycles. Its energy density on a weight

basis is 50% greater (32.6 Wh/kg). The relatively low operating potential of the MnO2 positive means

that it operates in a mild environemt from the viewpoint of corrosion, ensuring long life with graphite

as an electronically conducting element in the structure. Microcalorimeter studies have demonstrated a

low self-discharge rate for the positive. In view of the above, and taking into account the difficulty of

making nickel positives that can reliably offer 50,000 or more cycles, the secondary MnO2-H2 system

deserves serious consideration for LEO use.

REFERENCES

Klein, M.i Costa_ R.L." "Electrolytic Regenerative H2-O2 Secondary Fuel Cells,", Proc. Space

Technology and Heat Transfer Conf., Los Angeles, CA, 1970, American Society of Mechanical

Engineers, New York, 1970, 70-Av/SpT-39.

2. Appleby, A.J.: J. Power Sources 22, 377 (1988).

3. Klein, M.: Proc. 7th IECEC, Nat. Inst. Chem. Eng., Washington, D.C. 1972, p.729017.

4. Tsenter, B.I.; Sergeev, V.M.; Kloss, A.I.: U.S. Patent 3,669,744 (1972).

5. Dunlop, J.D.; Van Ommering, G.; Earl, M.W.: in D. H. Collins, ed., "Power Sources 6,"

Academic Press, London, 1977, p. 231.

6. Appleby, A.J.; Crepy, G; Feuillade, G.: Final Report, INTELSAT Contract No. IS-611,

Comsat Laboratories, Clarksburg, MD, 1975.

7. Beauchamp, R.: U.S. Patent 3,653,067 (1972).

8. Lim, H.S.; Verzwyvelt, S.A.: J.Power Sources 22,213 (1988).

o Kordesch, K.; Gsellman, J.; Peri, M.; Tomantschger, K; Chemelli, R.: Electrochim. Acta 26,

1495 (1981).

64

z
10. Yao, Y.J.; Gupta, N; Wroblowa, H.S.: J. ElectroanaI. Chem. 223, 107 (1987).

11. Wroblowa, H.S.; Gupta, N.: J. Electroanal. Chem. 238, 93 (1987).

12. Dziecinch, M.A.; Gupta, N.; Wroblowa, H.S.: J. Electrochem. Soc. 135, 2416 (1988).

65
TABLE 1: CELL WEIGHT BREAKDOWN IN GRAMS

Component Nominal Ni-H 2 LEO Ni-H 2 LEO MnO2-H 2

Positives

Hydrogen Electrodes

Separators

Gas Screens

Electrolyte

Stack Total

Center Rod and Insulator

Busbars and Tabs

Endplates

Nuts, Washers

Terminal Conductors

Internal Hardware

Pressure Shell

Weld Ring

Compression Seals

Container

Total Cell Weight

=- ;

Total Cell Voiume (ml)

348.6

72.0

35.1

8.0

159.0

622.7

14.3

41.4

34.0

4.6

79.0

173.3

179.0

35.0

18.4

232.4

1028.4

836.0

870.5

180.0

87.8

20.0

397.5

1555.8

35.8

41.4

34.0

4.6

79.0

194.8

275.4

35.0

18.4

328.8

2079.4

1286.0

154.4

72.0

35.1

8.0

159.0

428.5

14.3

41.4

34.0

4.6

79.0

173.3

170.5

35.0

18.4

223.9

825.7

805.0

Specific Energy (WhNg)

Specific Energy (Wh/l)

44.9

51.0

22.2

33.1

32.6

33.5

66
IOO

CYCLE NUMBER

Figure l: Reactant utilization versus cycle number for a modified MnO2 electrode cycled at 6C in 9M

KOH at room temperature

0 10 20

DISCHARGE TIME (HOURS)

Figure 2: First discharge curve for electrolytic MnO2 and modified MnO2 obtained in 9M KOH at

room temperature with a discharge current of 44 mA/gMn. Cathode mix: 85 wt%

manganese dioxide, 10 wt% acetylene black and 5 wt% binder.

67
Oil)

E

:f

o_

<

iv"

iii

Z

ILl

I--

<

W

7"

11.

0

UJ

l--

n"

120

8O

4O

31% KOH

25°C

MODIFIED-MnO2____.o

0 _ I i I I I

0 100 200 300 400

HYDROGEN PRESSURE, psi

Figure 3: Heat generation rates as a function of hydrogen pressure from the reactions between

hydrogen and NiOOH and bismuth-modified MnO2. 0.1g oxide samples in 1.0 ml of 31

wt% KOH Were used.

z

E

|

=

=_

L

68

m
N90-20463

CHARACTERIZATION TESTING OF A 40 AHR BIPOLAR

NICKEL-HYDROGEN BATTERY

Jeffrey C. Brewer

NASA Marshall Space Flight Center

Huntsville, Alabama 35812

Michelle A. Manzo

NASA Lewis Research Center

Cleveland, Ohio 44135

Russel P. Gemeiner

NASA Lewis Research Center

Cleveland, Ohio 44135

Extensive characterization testing has been done on a second 40 amp-hour (Ahr),

10-cell bipolar nickel-hydrogen (Ni-H2) battery to study the effects of such

operating parameters as charge and discharge rates, temperature, and pressure,

on capacity, Ahr and watt-hour (Whr) efficiencies, end-of-charge (EOC) and midpoint discharge voltages. Testing to date has produced many interesting results,

with the battery performing well throughout all of the test matrix except during

the high-rate (5C & IOC) discharges, where poorer than expected results were

observed. The exact cause of this poor performance is, as yet, unknown. Small

scale 2" x 2" battery tests are to be used in studying this problem.

Low earth orbit (LEO) cycle life testing at a 40% depth of discharge (DOD) and

I0°C is scheduled to follow the characterization testing.

INTRODUCTION

Space power systems of the future are projected to require power levels that

extend far beyond the current levels of demand. In order to meet these

increasing needs, improvements must be made to current energy-producing systems

or new technologies must be developed. Over the past several years, LeRC has

been actively engaged in the development of a bipolar configured Ni-H 2 battery.

This battery system has the potential to meet some of these high-power needs of

the future. In a continuing effort to develop this technology to a point where

it can be used efficiently in space flight, LeRC has begun testing a second 40

Ahr, 10-cell bipolar Ni-H 2 battery.

Results from the tests on the first battery tested here at LeRC were very

encouraging. The battery operated for some 10,000 LEO cycles at 40% DOD and

produced promising results in most of a variety of characterization tests, as

well (ref. 1). Following the completion of this test, work began on the design

of the second bipolar battery in hopes of developing an improved battery.

69
BATTERYDESIGN

The basic design of this battery differs from that of the first battery only

through slight modifications in the cell frames. Poor high-rate discharge

performance and electrolyte leakage paths in the first battery led to these

changes. It was thought, at the time, that a possible cause of the poor highrate performance in the first battery was limited gas access to the reaction

sites on the negative electrode. Gas-access area was thus increased by modifying

the cell frame in an attempt to alleviate this problem. The gaskets were also

changed in the cell frame design in an effort to improve the seals and thus

better contain the electrolyte within the cells, in addition to these minor

design changes, most of the individual components of the second battery came from

different manufacturers (table I). The reasons vary as to why these component

changes were made. The frame material, for instance, was changed from

polysulfone to ABS (resins which are terpolymers of acrylonitrile, butadiene,

and styrene) because of superior machining capabilities and mechanical stability

of the material. The nickel electrode was changed simply because of

availability. The other changes were based, at least in part, on more technical

reasoning and on the desire to reduce the number of parts. Because of

inconsistencies found in the previous H_ electrode, Giner, Inc. was engaged in

a development program to manufacture a sultable large area, single unit electrode

for this application. In doing so, the previous three-piece electrode would be

eliminated. A program was also undertaken with National Standard to develop a

fibrex electrolyte reservoir plate (ERP) which would contain pores in the desired

range and could be manufactured in one piece. The previous ERP was constructed

from nickel foam from Brunswick wl_ic_, due to the large area required and

manufacturing limitations, resulted in a six-piece ERP. In an attempt to

increase the effective current carrying area between the gas screen and the

bipolar plate and to improve high rate performance, the gas screen was changed

to a heavier, woven screen, as opposed to an expanded metal (Exmet) screen in

the first battery. This change also created a large weight increase which makes

it difficult to justify its use without major perf6rman_e improvement_. Finaiiy,

the electrolyte concentration was changed from 31% to 26% potassium hydroxide.

This was done because of superior life seen in IPV Ni-H_ testing (ref. 2). These

multiple component changes, as well as different testing procedures and unique

cell characteristics make it difficult to directly compare results from the two

batteries. Some comparisons, however, are valid and will be made.

One feature consistent with the first battery is that both utilized an active

cooling process. This is accomplished by pumping a coolant through alternate

bipolar plates (cooling plates) in the battery stack. This enables temperature

readings to be controlled very consistently and accurately throughout the entire

cell. This is one advantage over IPVtec_nology. In an IPV cell, temperatures

in the stack can run 7°C hotter than the measured temperatures outside the cell

(ref. 3).

PROCEDURES

After construction was complete, the battery was placed in a boiler plate

pressure vessel, which was designed to meet safety requirements. Each cell was

!

70

E
instrumented with voltage leads to measure individual cell voltages as well as

a thermocouple attached to the bipolar plate in the hydrogen cavity to measure

internal cell temperature. Additional thermocouples were placed on the exterior

of the stack to get additional temperature measurements. The pressure vessel

was equipped with a pressure transducer for measuring the hydrogen pressure, an

oxygen sensor, and a relative humidity indicator. A Modicon programmable

controller was used to run the test. All instrumented points were scanned and

digitized every IB seconds by a Neff multiplexer and stored by a central computer

system for subsequent processing. The data could then be processed and received

in both tabular and graphical form for each cycle.

Following the setup and checkout of the test hardware and data collection system,

initial cycles were run to determine battery capacity. An initial capacity of

40 Ahrs was assumed. The capacity determination cycles basically consisted of

a C/2 rate charge with a 5% overcharge followed by a C rate discharge to a

battery voltage of 7.0 V or a low cell voltage of 0.5 V. "C" is defined as the

rate at which the battery's capacity will be depleted in one hour. The low cell

voltage cut-off of 0.5 V was used in order to prevent any cell from going

negative and thus generating hydrogen. A C/4 drain to these same cut-off points

followed this to complete the total cycle. As the cycles continued, the "C"

value was adjusted several times until a consistent capacity was recorded, which

was 40 Ahrs. While some of the formation cycles were run at 20°C and 200 psi,

the baseline capacity determination cycles were run at ]0°C and 200 psi.

Characterization tests were then run at a variety of charge and discharge rates,

temperatures, and pressures. A full set of tests was run at ]0°C and 200 psi

at charge rates of C/4, C/2, and C, and discharge rates of C/4, C/2, C, 2C, 5C,

I0C, and 5C pulse for each charge rate. Following this set of tests, subsets

of this base characterization test were run at other temperatures and pressures.

A C/2 rate charge was chosen as the charge rate to be used in these subsets.

This selection was based not only on performance in the base characterization

test, but also on performance in prior tests (ref. 4). A C/2 rate charge also

allowed the cycles to fit better into an eight-hour day than if run at a lower

rate. A subset of tests run at 20°C and 200 psi consisted of the identical

discharge rates used in the base set paired with the C/2 charge rate. The

remaining subsets (O°C at 200 psi, 30% at 200 psi, and 20°C at 400 psi)

consisted only of the C/4, C/2, C, and 2C rate discharges. Again, all were

paired with the C/2 charge rate. Each individual test was run until three

consistent cycles were recorded. Consistency was based on Ahr and Whr

efficiencies. Each cycle consisted of a full charge (the amount of which was

equal to the capacity out in the preceding discharge plus a set percentage of

overcharge) followed by a full discharge to a battery voltage of 7.0 V or a low

cell voltage of 0.5 V (hereafter known as the normal cut-off points). A C/4

drain followed all discharges run at a C/2 rate or higher.

A set percentage of overcharge was used in this test in order to ensure adequate

charging as well as protect from unnecessary overcharging. This is compared to

the first battery test where a set charge input was used for each cycle

regardless of the capacity delivered in the previous discharge (ref. I). A 5%

overcharge was used initially; however, this proved to be insufficient to

adequately recharge the battery so the overcharge was increased to 10%. This

percentage maintained a stable capacity from cycle to cycle; however, to reduce

71
any unnecessary overcharging, the percentage was dropped to 7% where stable

capacities were once again realized. The overcharge amount thus stayed at 7%

throughout the remainder of the characterization testing.

Due to a lack of adequate manpower, cycles were not run over the weekends or

holidays. This extended wet-stand period allowed time for a weak, highresistance cell to self-discharge considerably more than the other cells. This

delayed the resuming of the characterization testing until after that cell could

be brought back up to a state of charge similar to that of the other cells. It

was found through trial and error that the weak cell could be maintained by

trickle charging the fully charged battery at aC/iSO to C/]O0 rate. This method

produced much more consistent results than other methods that were tried and

allowed characterization cycles to resume much more quickly.

RESULTS AND DISCUSSION

Due to the wide range of variables in this test, all of the relevant numerical

data will not be enumerated here, but is summarized in tables II - VIII. Figures

I - 4 contain pertinent voltage profiles from throughout the test.

Before discussion can begin on the specific results in the test, a general po!nt

needs to be made concerning the data to be discussed hereafter. It was mentiofied

earlier about the problems caused by a weak cell during extended wet-stand

periods. This same weak cell created problems during the characterization

cycles, as well. The discharges on most, if not all, of the early

characterization cycles were terminated by a battery voltage of 7.0 V. This

allowed good, accurate comparisons of data from cycle to cycle. The discharges,

however, on the vast majority of characterization cycles were terminated by a

low cell voltage of 0.5 V, while the overall battery voltage ranged from 7.1 to

10.4 V. Because most of the cycles thus had no common end-of-discharge (EOD)

point, it was difficult to compare the basic, overall data between cycles. So,

where it was helpful, capacity delivered to the 10.0 V point in the discharges

was used to compare cycles in hopes of negating some of the distorting effects

of the weak cell on the normal cut-off point data. Just what caused this cell

to perform this way is unknown at this time; however, it is thought that shunt

currents could be present which allowed an additional discharge path through

which the cell self-discharged over night between cycles. This would explain

the erratic behavior seen throughout the characterization cycles.

One additional comment about the data -- each data point represents the average

value of the three most consistent cycles run at that particular set of test

conditions.

Increasing the charge rate had little consistent effect on the capacity delivered

to the normal cut-off points, although at lower charge rates the battery seemed

to perform slightly better. Even the capacity delivered to I0.0 V shQwed no

consistent trends (table II). Increasing the charge rate also had little

consistent effect on the Ahr efficiency, but, due to the accompanying increase

in battery charge voltage from an average of 15.52 V at the C/4 rate to 16.65

V at the C rate, caused an average decrease of 7.3% in the Whr efficiency, except

72

L
at the 2C rate discharge where no significant changes were seen (table Ill).

Thus, because of its desirable effect on Whr efficiency and lack of effect on

other variables, a low charge rate would seem to be best; however, previous

testing has suggested that a C/2 to C charge rate range should produce the

optimum results (ref. 4). It is difficult, however, to directly compare results

from this characterization test with the referenced test due to the fact that

the nickel electrodes used in both tests came from different manufacturers.

Electrolyte concentration was also different. These differences alone could

account for the discrepancies seen between the two tests. In either case,

though, results failed to show strong proof that there is clearly an optimum

charge rate.

Increasing the discharge rate caused a consistent decrease in the capacity

delivered over all temperatures and pressures except at 20°C and 400 psi where

the capacity increased slightly from 46.03 Ahr to 47.46 Ahr at discharge rates

of C/4 and C/2, respectively (table IV). The probable cause of this apparent

increase in capacity delivered at the higher rate can be traced to the weak cell

causing an early termination of the discharges run at the lower rate. Comparing

the capacities at the 10.0 V point in the discharge supports this theory. At

10.0 V, an average capacity of 43.gg Ahr was delivered at the C/4 rate and 42.59

Ahr was delivered at the C/2 rate.

Because of the lower capacities delivered at the higher discharge rates, the

resulting Ahr efficiencies also decreased across the board (table V). Also, as

discharge rates increased, Whr efficiencies decreased over all temperatures and

pressures due to the decrease in operating voltages that always accompany

increasing discharge rates (tables VI &VII). In this report, operating voltages

are reported as mid-point discharge voltages, which were calculated by averaging

the following two data points: the voltage reading at I/2 of the total discharge

time and the voltage reading at the 20 Ahr out point. Discharge voltage profiles

vs. capacity at all discharge rates can be seen in figure 1.

The discharge rate effects were all as anticipated; however, at high discharge

rates of 5C and 10C, performance was very poor. Poor performance was also seen

at these rates in the first battery (ref. 1); but, despite attempts to alleviate

this problem through design and individual component changes, even poorer

performance was seen in the second battery (figure 2). In the attempts to run

a 10C discharge, the battery voltage dropped below 7.0 V within 30 seconds.

Another 10C discharge was run and was allowed to continue past the normal cutoff points down to a low cell voltage of 0.1V. This discharge lasted 3 minutes

but the voltage did not begin to level off until around 4.5 V (figure 1). The

5C discharges lasted longer but, again, failed to level off above a battery

voltage of 7.0 V. (tables IV - VI and figure 1). Only a 5C pulse (1 second on

/ 4 seconds off) discharge was able to produce meaningful results (tables IV &

V and figures ] & 3).

Several ideas have been discussed as to what could be causing this high-rate

discharge problem. Limited gas access to the negative electrode was previously

mentioned as a possible cause. The design changes mentioned earlier that were

made in an attempt to alleviate this problem instead could have elevated the

problem even more. This is based on the possibility that the holes drilled in

the battery frame to allow gas access into the interior of the battery became

73
filled with electrolyte and blocked free gas flow. One argument against this

scenario is that the pressure of the gas flowing to and from the electrodes would

keep the holes clear. Also, if gas access is the problem, then, based on the

amount of H_ located in the cavity adjacent to the negative electrode, initial

voltage performance might be expected to be good but would fall off as the H2 gas

located in the cavity was used up. Data from the high-rate discharges showed

no signs of good initial voltage performance and even showed signs of leveling

off at a low voltage. A second possible cause is poor contact between the gas

screen and the two surrounding components -- the negative electrode and the

bipolar plate. Due to the large surface area of the Bipolar stack compo-nents_

uniform stack compression and the resulting surface contact between individual

components is difficult to consistently maintain. A lack of adequate contact

area would limit the current carrying capability and produce poor results

especially at high current levels. Another possible cause is the Goretex backing

placed on the H. electrode during the standard production orocess at Giner Inc L , . , •

Its presence could posslbly be limiting the effective Contact area between the

H2 electrode and the gas screen, thus hindering current flow. There was some

initial concern about the possibility that this backing could cause current flow

problems in a bipolar configuration; however, polarizations of up to 500 mA/cm 2

were operated successfully on a small scale prior to construction of the full

size electrode. Perhaps the key point, once again, is the possi6ie lack of

uniform stack compression in a full scale battery configuration. One possible

solution that has been discussed that could, at least, partially improve the

contact between components is to weld the gas screen to the bipolar plate. This

would assure adequate contact area between these two components, but Would not

improve the contact area between the H_ electrode and the gas screen. The

effects on performance that the Goretex backing has in a full scale battery as

well as other possible problem areas are to be addressed in further testing.

Small scale 2" x 2" battery tests will begin soon and will be used to evaluate

some of these areas. Flooded capacity tests have already been done to evaluate

the performance effects caused by varying the nickel electrode manufacturer and

the electrolyte concentration. Electrodes from both manufacturers (Eagle Picher

and Whittaker-Yardney) were tested at ]OC, SC, 2C, C, C/2, and C/4 discharge

rates using 26, 31, and 40% KOH as electrolyte. Nickel was used as the counter

electrode and amalgamated zinc was used as the reference electrode. Results from

this test showed no signs of inferior high-rate performance by the Whittaker'

Yardney nickel electrode used in the second battery (figure 4). On the contrary,

these electrodes produced much more stable efficiencies and capacities at all

discharge rates tested (table VIII). The Whittaker-Yardney electrodes shown in

table VIII produced 76% of their low-rate (C/_capacity atthe high (IOC) rate,

while the Eagle Picher electrodes delivered only 39% of their low-rate capacity

at the high rate. It is not possible, however, to completely rule out the nickel

electrode as being responsible for the poor high-rate discharge performance,

Because of the flooded conditions under which these capacity measurements Were

made, the ability of the different electrodes to perform under actual battery

conditions was not addressed. Thus, it is entirely possible that under actual

battery conditions the Whittaker-Yardney electrodes would not perform optimally

and that the "starved" condition could lead to the type of poor performance that

was seen at the high rates. Finally, it was not intended through these tests

to directly compare the two manufacturers' electrodes. Neither manufacturer

optimized the electrodes that were used; they simply supplied standard electrodes

of the size requested.

74
Temperature variations between O°C and 30% produced some interesting results.

As expected, an increase in temperature produced significant drops in EOC voltage

due to the decrease in internal resistance that accompanies rising temperatures.

Readings averaged 16.40 V at O°C and 15.12 V at 30°C. As temperatures increased,

the mid-point discharge voltages, however, showed a steady increase at all

discharge rates except at C/4 where a 50 mV drop was seen between the

temperatures of O°C and ]O°C. (table VII). The improved voltage performance seen

at higher temperatures, however, did not directly translate into an increase in

Whr efficiencies at all conditions (table VI). At the 2C rate discharge the Whr

efficiency was the lowest (41.26%) at 30°C and the highest (53.43%) at 20°C. At

the other three discharge rates, though, 30°C produced the highest efficiencies.

The discrepancy at the 2C rate discharge can be, at least partially, attributed

to a high EOD voltage of ]0.3 V at 30% compared to 7.2 V at 20°C. Whether or

not the large percentage difference could have been completely overcome had the

30% discharge run down to around 7.0 V is difficult to determine; but, certainly

a large portion of it would have been. The effect that temperature had on total

capacity out was equally interesting. For instance, as temperatures increased

from O°C to 20°C, capacity delivered to the normal cut-off points increased at

the C/2 rate discharge from 41.]9 Ahr to 47.68 Ahr, respectively, but fell off

drastically at 30°C to 37.50 Ahr (table IV). This trend was consistent at all

discharge rates and was also seen in the capacity data to the 10.0 V point in

the discharge. Ahr efficiencies, however, seemed to be less consistently

affected by temperature variations (table V). All of this data seems to support

the use of temperatures as high as 20°C or even 30°C to produce optimum results.

This agrees rather well with the data produced during the first battery test

(ref. I).

Increasing the H_ pressure inside the vessel also produced some interesting

results. Negllgi_le change was seen in the EOC voltage between 200 and 400 psi

as the voltage dropped from 15.62 V to ]5.59 V, respectively; however, at 400

psi, an improvement was seen in the mid-point discharge voltage at all discharge

rates (table VII). This was expected behavior because the increased pressure

would increase the activity coefficient of the gas and thus improve its

efficiency and voltage performance. When looking at the data measured to the

normal cut-off points in the discharge, the capacity, and Ahr and Whr

efficiencies all were less at the higher pressure, except at the C/4 rate

discharge, where a slight increase in both Ahr and Whr efficiencies was seen

(tables IV - VI). These were not expected results but, once again, the weak

cell seems to be distorting the data by prematurely terminating the discharges

during the 400 psi cycles. Although the capacity delivered to ]0.0 V is still

greater at 200 psi, the differences are not as great. Also, the Ahr and Whr

efficiency differences can be reasonably eliminated by considering the high EOD

voltages on the 400 psi cycles. Actually, the Whr efficiencies would have

probably been greater at 400 psi had all discharges terminated at similar

voltages. It certainly should illustrate that increasing or decreasing the

pressure will have minimal effects on overall battery performance.

After completion of the characterization cycles, the battery was placed on LEO

cycle life testing at 40% DOD and at ]O°C.

75
CONCLUSIONSANDRECOMMENDATIONS

Despite the fact that a weak cell made it difficult to directly compare some of

the data from cycle to cycle, there was enough evidence to see that the battery

produced generally expected results and performed very well throughout the

majority of the characterization test matrix. It is hoped and expected that the

LEO cycle life test that has just begun will produce _imilarly encouraging

results. One area that continues to be a problem, however, is the high-rate

discharge performance of the battery. Even though improvements were not made

in this area with this battery, several encouraging ideas have been mentioned

as possible solutions to the poor high-rate performance problem. As mentioned

earlier, several studies, including small scale 2" x 2" battery tests, will be

done in hopes of pinpointing the area or areas responsible for the poor high_

rate performance.

REFERENCES _:_:__ _

1. Cataido,_R.F.-_-r_metric and Cycle Tests of a 40'A-hr B_o_ar Nickel-Rydrogen

Battery. 21st Intersociety Energy Conversion Engineering Conference, Vol.

3, pp. 1547-1553.

2. Lim, H.S.; and Verzwyvelt, S.A.: KOH Concentration Effect on the Cycle Life

of Nickel-Hydrogen Cells. II. Accelerated Cycle Life Test. 21St Intersociety

Energy Conversion Engineering Conference, Vol. 3, pp. 160]-1606.

3. Rogers, H.; Levy, Jr., E.; Stadnick, S.J.; and et al: Failure Mechanisms in

Nickel-Hydrogen Cells, Final Report. December 1977, pg. 2-36.

4. Cataldo, R.L.: Test Results of a Ten Cell Bipolar Nickel-Hydrogen Battery.

18th Intersociety Energy Conversion Engineering Conference, Vol. 4, pp.

1561-1567.

76
Table I. COMPONENT DIFFERENCES BETWEEN BATTERY I AND BATTERY 2

Component Battery ] Battery 2

Frame Polysulfone - Old ABS New

Nickel electrode Eagle Picher Whittaker - Yardney

Hydrogen electrode LSI Giner

ERP Ni foam - Brunswick Ni felt Nat'l Std

Gas screen Exmet Woven (Nat'l Std)

Separator Asbestos Asbestos

Electrolyte 31% KOH 26% KOH

Table II. - CHARGE RATE EFFECTS ON CAPACITIES DELIVERED TO

NORMAL CUT-OFF VOLTAGES & TO 10.0 VOLTS AT I0°C

AND 200 PSI; BATTERY EOD VOLTAGES SHOWN UNDER

NORMAL CUT-OFF CAPACITY VALUES

Discharge

Rate

c14

c12

C

2C

Capacity to

Normal Cut-offs

Capacity to 10.0 V

Charge rate Charge rate

c/4 C/2 C c

46.90 46.87

6.9 7.0

46. ] 2 43.72

6.8 7.1

42.38 40.84

7.3 8.0

37.93 37.17

7.1 7.4

c/4 c/2

43.15 42.85

42.60 42.02

37.75 39.08

20.88 2].52

47.45

7.0

46.80

7.0

41.73

7.1

35.92

7.1

43.06

40.34

38.97

24.27

7?
Table III.- CHARGE RATE EFFECTS ON AHR & WHR EFFICIENCIES AT

]O°C AND 200 PSI

Discharge Ahr Efficiency, % Whr Efficiency, %

Rate

Charge rate Charge rate

c14 c12 c14 I c12 c

l

c14

c/2

C

2C

Table IV. -

94.16

87.76

76.99

67.09

93.60

86.30

77.04

67.72

C

93.77

85.95

75.76

70.39

80.16

72.43

60.43

46.62

78.]7

69.21

58.93

45.85

74.45

66.20

55.83

46.32

CAPACITY DELIVERED TO NORMAL CUT-OFF VOLTAGES AT ALL

TESTED CONDITIONS WITH A C/2 RATE CHARGE; BATTERY EOD

VOLTAGES SHOWN UNDER CAPACITY VALUES

Discharge Capacity, Ahr

Rate

200 psi 20°C

O°C I0°C 20°C 30°C 200 psi 400 psi

c14

c12

C

2C

5C

5C Pulse

42.15

9.7

41.]9

10.4

38.32

9.3

33.17

7.6

###

###

46.90

6.9

46.12

6.8

42.38

7.3

37.93

7.1

]5.43

7.0

30.60

7.0

50.49

7.4

47.68

***

44.20

7.3

39.45

7.2

22.24

6.9

33.97

8.1

40.7]

7.9

37.50

7.3

35.48

9.2

32.60

]0.3

###

###

50.49

7.4

47.68

44.20

7.3

39.45

7.2

22.24

6.9

33.97

8.1

46.03

8.9

47.46

7.5

38.82

8.7

33.66

g.o

###

###

=

=

readings available due to

5C and 5C Pulse run only

78

data collection system errors

at 10°C and 20°C at 200 psi
Table V. AHREFFICIENCIESAT ALL TESTEDCONDITIONSWITHA C/2

RATECHARGE

Discharge Ahr Efficiency, %

Rate

200 psi 20°C

O°C 20°C 30°C 200 psi 400 psi

c14

c/2

C

2C

5C

5C Pulse

90.80

85.79

82.80

72.67

###

###

lO°C

93.60

86.30

77.04

67.72

28.18

52.08

92.04

86.76

78.86

72.39

41.52

65.18

90.96

86.68

79.62

72.23

###

###

92.04

86.76

78.86

72.39

41.52

65.18

92.55

86.20

76.85

68.38

###

###

### - 5C and 5C Pulse run only at ]O°C and 20°C at 200 psi

Table VI. - WHR EFFICIENCIES AT ALL TESTED CONDITIONS WITH A C/2

RATE CHARGE

Discharge Whr Efficiency, %

Rate

200 psi 20°C

O°C I0°C 20°C 30°C 200 psi 400 psi

C/4

C/2

C

2C

5C

75.90

69.28

61.67

46.51

###

78.]7

69.21

58.93

45.85

15.35

77.92

63.69

53.43

24.82

79.99

75.29

68.32

41.26

###

77.92

***

63.69

53.43

24.82

80.09

71.52

63.64

52.91

###

*** No readings available due to data collection system errors

### - 5C and 5C Pulse run only at I0°C and 20°C at 200 psi

79
Table VII. - MID-PT. DISCHARGE VOLTAGES AT ALL TESTED CONDITIONS

WITH A C/2 RATE CHARGE

Discharge Mid-Pt. Discharge Voltage, V

Rate

200 psi 20°C

k,

O°C 10°C 20°C 30°C 200 psi 400 psi

C14

C/2

C

2C

12.90

32.40

12.85

12.45

12.95 13.20

12.65 13.00

12.95

12.65

11.45

9.70

11.72

]0.15

12.15

11.10

12.60

11.83

12.15

11.10

13.10

12.80

12.45

11.55

Table VIII. - CAPACITIES & AHR EFFICIENCIES FOR FLOODED ELECTRODE

TESTS ....

Discharge Capacity, Ahr Ahr Efficiency, %

Rate

EP31 WY26 EP31 WY26

c/4

C/2

C

2C

5C

10C

0.64

0.63

0.59

0.55

0.47

0.25

0.43

0.41

0.40

0.39

0.37

0.33

71.95

71.19

66.54

61.92

52.63

28.20

73.07

70.48

68.00

67.43

62.86

56.90

80

B

h
v

16 -

ZO t0

\\

Cepac£ty (_r)

Figure 1. - Discharge voltage profiles vs. capacity for all

discharge rates @ 10% & 200 psi

16

15

13

12

OC/! "_'_.

8

• oc/,2

"_ -.-.....:_..._

5C/#2 -m

Figure 2.

10 7.0 3O

_p|cltY ( ,_lr )

5C l IOC rate discharge voltage profile vs. capacity @

20% & 200 psi for battery 1 & battery 2

8]
i6

14

13

== lz

_0

7

13 l O 20 :30

CIptcLt? (k_r)

Figure 3. - 5C pulse (1 sec on / 4 sec off) discharge voltage

profile vs. capacity @ 10°C & 200 psi

3

j,.

2

_..9

L.8

L,7

_.6

l..t5

f-.4

1,3

1.2

£.t

.l

0.9

O.e

0.7

0.6

0.5

0.4

0.3

.,.-,_ _-_ _.

_--.,._. "k

- vYze/loc

\

0 0.2 0,4

Clptcl.t,y ( ).lie )

Figure 4. - IOC & 5C d|s_a_e Voltage profiles vs. capacity

during flooded electrode tests for Eagle Picher

electrode in 31% KOH and Whittaker/Yardney electrode in 26% KOH

82
N90-20464

ADVANCED DOUBLE LAYER CAPACITORS

S. Sarangapani, P. Lessner, J. Forchione,

A. Griffith and A.B. LaConti

Giner, Inc.

Waltham, Massachusetts 02254

Work has been conducted that could lead to a high energy density electrochemical capacitor, completely free of liquid electrolyte. A three-dimensional

RuOx-ionomer composite structure has been successfully formed and appears to

provide an ionomer ionic linkage throughout the composite structure. Capacitance values of approximately 0.6 F/cm z have been obtained compared I F/cm 2 when

a liquid electrolyte is used with the same configuration.

INTRODUCTION

Electrochemical capacitors store energy by utilizing double layer and surface redox type processes. Table I compares energy density, power density, and

cycle life figures for batteries, electrochemical capacitors, and conventional

capacitors. The projected energy density for electrochemical capacitors are two

orders of magnitude lower than that of batteries, but power densities (on the

millisecond to second) time scale are three orders of magnitude higher. Energy

density is much better than for conventional capacitors, but the nature of the

electrochemical processes make the electrochemical capacitors suitable for relatively long pulses (milliseconds) and low to intermediate power applications.

Table I

Device Wh/L" W/L Cycle Life

Batteries 50-250 150 <104

Electrochemical Capacitors 5 >105 >i05

Conventional Capacitors 0.05 >i08 >106

Several applications of devices with the characteristics of electrochemical

capacitors can be envisioned: lightweight electronic fuses, burst power for

battery-powered vehicles and backup power for computer memory. The last application takes advantage of the high cycle life of electrochemical capacitors,

instead of the high power density. Commercial devices, based on high surface

area carbons and sulfuric acid and marketed by NEC under the SUPERCAP brand

label, are already being sold for this application. Preliminary projections

83
indicate that the advanced electrochemical capacitors that are being developed

will have an order of magnitude higher specific energy and energy density than

these carbon-based capacitors.

A simplified equivalent circuit of a single electrochemical capacitor cell

is shown in Figure I. A device consists of two electrodes separated by an electrolyte. Capacitances for metals and carbons in contact with aqueous solutions

are 10-40 microfarads/real cm 2. Certain noble metal oxides in contact with

aqueous solutions have capacitances in the range of approximately 150 microfarads/cm 2 (Kleijn and Lyklema, 1987). Materials in contact with organic

solvent-electrolytes may also have high capacities. In addition, pseudocapaeity

due to fast surface redox reactions can contribute to energy storage which is

accessible on a similar time scale as double layer energy storage. High surface

area porous electrodes can be used to give capacitances up to several farads per

geometric cm 2.

The equivalent circuit also illustrates several resistances present in the

system. The resistance in parallel with the capacitance represents a leakage

path. It should be as high as possible. The series resistance between the

electrodes ultimately limits the power that can be delivered by the system.

The major contributor to this is the resistance of the electrolyte between the

electrodes. In addition, for practical systems using porous electrodes, the

distributed electrolyte resistance in the electrode contributes to the series

resistance,

Because of the decomposition of water, aqueous-based devices are limited to

about 1 V per cell; organic-based devices may have higher cell voltages because

organic solvents have higher decomposition voltages. To build up the voltage,

several cells need to be stacked in series.

These physical principles provide an explanation for some of the figures of

merit listed in Table I. Energy density of electrochemical capacitors is lower

than that of batteries because the capacitors use surface processes to store

energy instead of bulk processes. However, these surface processes are accessible on a much shorter time scale than bulk processes giving the electrochemical

capacitor a higher power density. Cycle life is greater than that of batteries

because the surface processes result in no or small morphological changes in the

material. In a sense, the electrochemical capacitor can be thought of as a

specialized type of battery, r

Some work has been conducted in the development of solvent-electrolytebased electrochemical capacitors (Holleck,et al., 1988; Lee, et al., 1988).

A problem with these systems is the possibility of electrolyte leakage. This

can occur because of external pressure changes or gas evolution due to ceil

overcharge. The typical sulfuric acid electrolyte (3 to 4M H2SO4) is very corrosive and leakage could cause the module and ancillary equipment to be damaged

or destroyed. This has led us to investigate devices where the liquid electrolyte is replaced by a solid ionomer electrolyte. Perfluorosulfonated ionomer

electrolytes, such as Nafion (R), consist of a perfluoro polymer backbone to

which sulfonlc acid sites are permanently anchored. The only liquid or vapor

present in these devices would be pure water.

(R) E.I. DuPont Trademark

84

e

Z
The device concept is shown in Figure 2. The capacitor electrodes (high

surface area metal oxides) are separated by thin sheets of ionomer membrane.

Current collection is accomplished via lightweight foils or screens. The configuration allows the individual cells to be stacked into a module. The repeating element consists of the thin bipolar metal collector in intimate contact

with a unitized particulate-solid ionomer composite structure.

The major challenge to be met when replacing the liquid electrolyte with

the solid ionomer electrolyte is maintaining a high surface area electrodeelectrolyte contact. A liquid electrolyte is sorbed into the macropores of the

electrode and mlcropores down to about i00 A by capillary action. An illustration of the type of composite that needs to be formed with the solid electrodesolid ionomer electrolyte structure is shown in Figure 3. A film of ionomer is

needed from the face of the electrode in contact with the ionomer membrane

separator extending back into the electrode toward the current collector. The

electrode particulates must make intimate electronic contact with each other and

with the current collector. Ideally, there is ionomer ionic linkage throughout

the thin particulate-solid ionomer composite structure and extends from metal

collector to metal collector.

There are several requirements for the electrode material to be used in the

electrochemical capacitor. Most importantly, it should have a high double layer

capacity and pseudocapaclty on a real area basis. In order to have a high

capacitance on a geometric area basis, the material should have a high surface

area. The material should be able to be fabricated into a porous electrode

forming a high surface area interface with the solid ionomer electrolyte and

bonded to the ionomer membrane separator.

RuO 2 was selected as the electrode material to be studied initially because

it meets all of these requirements. The oxide has a high double layer capacity

of about 150 microfarads/real cm 2 (Kleijn and Lyklema, 1987). There is also a

substantial pseudocapacity due to the surface reaction which can be written

approximately as:

2RuO 2 + 2H + + 2e- --> Ru203 + H20 [i]

This type of surface reaction in combination with double layer processes has

been shown to be capable of sustaining high current densities on the millisecond

scale (Holleck, et al., 1988; Lee, et al., 1988). RuO 2 can be made into a high

surface area coating (Trasatti and Lodi, 1980) or as high surface area particulates (Balko, et al., 1980). Bonding of particulates to an ionomer membrane has

been demonstrated for RuO 2 anodes for use in chlor-alkali cells (Coker, et al._

1980) where membrane and electrode assemblies with active areas of up to 35 ft =

have been prepared.

EXPERIMENTAL METHODS

RuO 2 was prepared by thermal decomposition of RuCI 3 in a NaNO 3 flux (Balko,

et al., 1980). The resulting powder was characterized by scanning electron

microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction

(XRD) and nitrogen surface area (BET method).

85
Two approaches were used to form the RuO2-1onomer composite electrodes:

i) impregnation of a PTFE-bonded RuO 2 electrode with ionomer solution, and

2) mixing of ionomer solution with RuO 2 powder, followed by formation into an

electrode. In all cases, the ionomer solution was 5% Nafion 117 in an alcoholwater mixture (Aldrich Chemicals).

The RuO2-Nafion composite electrodes were hot-pressed to a Nafion 117

membrane. A diagram of the test cell is shown in Figure 4. The electrode

arrangement for testing consisted of-two electrodes pressed on opposite sides

of the membrane. The membrane and electrode (M and E) assembly was held between

two acrylic piates which contained a Water reservoir to keep the M and E wet.

A tail of the Nafion 117 membrane dips into a beaker of H2SO 4 which contains a

Hg/Hg2SO 4 reference electrode. Some of the M and E assemblies had an integrally

bonded Pt/alr reference electrode.

The effectiveness of forming the extended RuO2-ionomer interface and the

maximum electrochemical limits of operation were assessed by using cyclic

voltammetry. T_he electrode was cycled between voltage limits where only double

layer and reversible surface redox processes occur. For double layer or surface

processes, the current is related to sweep rate by:

dV

i - C

dt

where C is the capacitance of the electrode.

[2]

RESULTS AND DISCUSSION

Powder X-ray diffraction showed the material to be substantially RuO 2.

There was some undecomposed _uCl 3 present. Nitrogen adsorption measurements

gave surface areas pear 50 m=/g. The SEM photograph in Figureb_s_0ws_the

particles to be pyramid-shaped and on the order of 0.i to 5 microns in size.

Clearly these surface areas cannot be explained on2the basis of impermeable

particles of this size. TEM examination of a 50 m /g powder (Figure 6) shows

that the particles have considerable internal porosity with crystallites on the

order of 200 A. Assuming spherical particles of this diameter, the calculated

specific surface area is about 40 m2/g, which is in good agreement with the

measured BET surfac_ _rea. _is powder was heat treate _ which reduced its

surface area to 3 m=/g. Figure 7 shows that this surface area loss was due to

the loss of the small size particles by sintering. This type _f morphology in

the oxide is not unexpected. The density of RuCI 3 is 3.1 g/cm _, while RuO 2 has

a density of 7 g/cm . When the decomposition reaction is accompanied by a large

decrease in specific volume (as is the case here) the final products tend to

have the external shape and dimensions of the precursor, but with many small

voids (Volpe and Boudart, 1985). This appears to be the case here.

Voltammograms of a RuO 2 electrode bonded to a Nafion 117 membrane are shown

in Figure 8. At potentials more negative than -0.4 V vs. Hg/HgS04, RuO 2 is

irreversibly reduced to lower oxides (shown as the beginning of a reduction wave

in Figure 8). Above about 0.5 V, oxygen evolution is possible. This gives a

single cell a usable voltage of 0.9 V. In the intermediate region, the voltammogram is fairly flat with peaks due to reversible surface conversion of oxides

86

z

m
as in Reaction i. Trassati and Lodi (1980) have stated that the oxide cycles

between the +3 and +4 states in the region of -0.3 to 0.3 V. The current in

the potential region near 0.15 V was plotted against sweep rate to obtain

capacitance.

The capacitances of 50 m2/g powder bon@ed to the Naflon 117 membrane with

no ionomer impregnation are about 0.15 F/cm . This capacitance is independent

of RuO 2 loading. This is approximately one order of magnitude lower than the

capacitance if the electrode was completely flooded with sulfuric acid.

Our most successful attempts to extend the area of electrode-electrolyte

contact has been via the technique of suspending RuO 2 in a Nafion solution and

spreading it onto the Naflon 117 membrane, followed by evaporation of the

solvent. Figure 9 shows the results obtained by using this technique with a low

surface area powder and 5 wt% Nafion in the final electrode. Although the absolute values of capacitances are low, a clear correlation is seen between RuO 2

loading and capacitance. Figure i0 shows that with a higher surface area

powder, similar results can be achieved. Another technique is to form the electrode on a separate thin sheet, followed by transfer to the membrane. Figure 10

shows that this method gives similar results; it _s also more amenable for

making larger electrodes. Electrodes up to 25 cm = have been made with this

technique.

SUMMARY

The all-solid electrochemical capacitor concept offers the advantage of

greater reliability and safety. If the electrode-ionomer structure can be

optimized further and sufficiently thin ionomer membrane separators can be

developed, then the energy and power density for the all-solid capacitor should

be comparable to other advanced electrochemical capacitors being developed (Lee,

et al., 1988 and Table I). The techniques of seallng/water management, thermal

management and power transfer/control are expected to be similar to scaled-up

all-solld-state proton exchange membrane bipolar electrolysis devices (C_ker, et

al., 1982). These devices containing up to I00 bipolar elements (0.3 ft =) have

been marketed as reliable devices for over 5 years.

ACKNOWLEDGEMENT

This work is being supported by the Office of Naval Research under Contract

No. N00014-88-C-0391.

REFERENCES

Balko, E.N.C.R. Davidson and A.B. LaContl. 1980. Solid Solutions of RuO 2 and

IrO 2. J. Inorg. Nucl. Chem. 42:1778.

Coker, T.G., R.M. Dempsey and A.B. LaConti. 1980. Production of Halogens in an

Electrolysis Cell with Catalytic Electrodes Binded to an Ion Transporting

Membrane and an Oxygen Depolarized Cathode. U.S. Patent No. 4,191,618.

87
Coker, T.C., A.B. LaConti and L.J. Nuttall. 1982. Industrial and Government

Applications of SPE (R) Fuel Cell and Electrolyzers. Paper published in

Proqeedings of the Electrochemical Society Symposium on Membranes and Ionic and

$_ectronic Conducting Polymers, Case Western Reserve, Cleveland Ohio.

Holleck, G.L., B.H. Jackman and R.D. Rauh. May 1988. High Rate Electrodes for

Pulse Power, Iridium Oxide and Ruthenium Oxide. Abstract No. 18, The 173rd

Electrochemical Society Meeting, Atlanta, Georgia.

Kleijn, J.M. and J. Lyklema. 1987. The Electrical Double Layer on Oxides:

Specific Adsorption of Chloride and Methylviologen on Ruthenium Dioxide.

Journal of Colloid and Interface Science. 120:511.

Lee, Ho-Lun, G. Mason and K. Kern. 1988. High Power Pulse Energy System

Concept. Final Report for ONR Contract No. N00014-87-C-0705.

Trasatti, S. and G. Lodi. 1980. Properties of Conductive Transition Metal

Oxides with Rutile-Type Structure. In Electrodes of Conductive Metallic Oxides

Fart A, Chapter 7, S. Trasatti (Ed.), Elsevier, Amsterdam.

Volpe, L. and M. Boudart. 1985. Topotactic Preparation of Powders with High

Specific Surface area. Catal. Rev. Sci. Eng. 27:515.

Figure I.

m

//I _--" "_

/ \

/ \\

I C(. Call t

-( I '--" II ,--

I |

t i I I

I

\1 't

I ,/

_\R Shunt R Shunl. /

Simplified Equivalent Circuit of Electrochemical

Capacitor Single Cell.

88
Ti FOIL HIGH SURFACE AREA OXIDE

SOLID ELECTROLYTE

Figure 2. Unitized Electrochemical Capacitor Module.

METAL OXIDE

PARTICLE

NAFION

TEFLON

PARTICLE

WIRE

Figure 3. Structure of lonomer-Metal Oxide Composite Electrode.

89
ORIGINAL PAGE iS

OF POOR QUALITY

GASKET

Pt REFERENCE ELECTRODE

CAPACITOR ELECTROOE

...... EEAD _:

...-.- I_ACKING PLATE

WATER

RESERVOIR

KET

IONOMER MEMBRANE

Figure 4. Test Fixture for All-Solld capacitor Cell.

Figure 5. SEM Photograph of RuO 2 Particulates.

90
ORfG|NAL P_GE _3

OF POOR QUALITy

m

Figure 6. TEM Photograph of 50 m2/g RuO 2 Particulates.

L

Figure 7. TEM Photograph of 3 m2/g RuO 2 Particulates.

9]
ORIGINAL PA_E !_

OF POOR QUAL!TY

I

\

\

I

\

\

i

o!

\ \

\

\

o o

VW '.l.N3_lanO

o

I

92
0

o.

I I I I

+

+

I t I I

0 O O 0

0 0 0 0

( L-3)

(_uJol.-l) eoutw31oodoo

-4-

0

o.

0

0

,If

140

W

0

t=

0

i::

0

+-

II

0

.J

411

0

O:

"U

0

L

(U

.e

r.-

O

z

i

-g

III

L

lU

U

"2

"-I

O

N

O

I.. tO IU

m E

C II;

O ,_

• m

tlJ a

u

e" O

I0 "O

ID O

(.

93
I I I I

_PJO

b.

+0

o

@4

co

@l

,IF=.

CO

R"

E

0

E

e"

mu

II

0

.,I

m

0

nW

ILl

0

L,.

U

0

U,,I

C

0

,4-

m

UlE

|

0

1.

ILl

U

e

N

0

-!

*v ,

(U

I. C

0 CO

C tU

a _

0 ,,-"

_J _

_ C

_ 0

U

C O

_ •

m _

n C

lid 0

I I 1 I 0

ILl

o o o o o o *-

o o o o 0 '_

(,WOld) eouu),!oudgo

94
N90-20465

RECENT ADVANCES IN SOLID POLYMER ELECTROLYTE FUEL CELL

TECHNOLOGY WITH LOW PLATINUM LOADING ELECTRODES

Supramaniam Srinivasan, David J. Manko, Hermann Koch 1,

Mohammad A. Enayetullah and A. John Appleby

Center for Electrochemical Systems and Hydrogen Research

Texas Engineering Experiment Station

Texas A&M University System

College Station, Texas 77843

Of all the fuel cell systems only alkaline and solid polymer electrotyte fuel cells are capable

of achieving high power densities (>lW/cm 2 ) required for terrestrial and extraterrestrial

applications. Electrode kinetic criteria for attaining such high power densities are discussed.

Attainment of high power densities in solid polymer electrolyte fuel cells has been demonstrated

earlier by different groups using high platinum loading electrodes (4 mg/cm 2) . Recent works at

Los Alamos National Laboratory and at Texas A&M University (TAMU) demonstrated similar

performance for solid polymer electrolyte fuel cells with ten times lower platinum loading (0.45

mg/cm 2 ) in the electrodes. Some of the results obtained at TAMU are discussed in terms of the

effects of type and thickness of membrane and of the methods of platinum localization in the

electrodes on the performance of a single cell.

1 INTRODUCTION

1.1 Rationale for Selection of Solid Pol mer E1 1 e Inst ad of Alkalin El I te F el

Cell System for Attainment of High Power Densities

Only the alkaline and solid polymer electrolyte fuel cell systems are capable of attaining

high power densities (>1 W/cm2). The pros and cons of the two fuel cell systems are best

expressed as in Table 1. Even though the alkaline fuel cell systems developed by International

Fuel Cells/United Technologies Corporation(l), for the Apollo and Space Shuttle Program, are in a

highly advanced state and have functioned extremely well for the required missions, the solid

polymer electrolyte fuel cell system is a stiff competitor for the alkaline fuel cell system, as is

clearly evident in Table 1. There is already a sufficient Technology Base Development for the

alkaline system at International Fuel Cells/United Technologies Corporation and several $100 M

have been invested in this system. Even though a considerably lesser financial support has been

provided for the solid polymer electrolyte system, it is in an advanced state of development, as

evidenced by the progress made in the development of 1.5 kW systems by Ballard Technologies,

Inc. in Vancouver, Canada (2) and by Ergenics Power Systems, Inc., in Wyckoff, New Jersey(3).

Siemens(4) in Erlangen, Germany is also engaged in research, development and demonstration of

high power density solid polymer electrolyte fuel cell systems for submarine applications. The

work on solid polymer electrolyte fuel cell systems by these organizations has been with high

platinum loading (4 mg/cm 2) electrodes. In the most recent work at Ballard, it has been reported

that a current density (i) of 6 A/cm 2 has been attained at cell potential (E) of 0.5 V. The slope of

the linear region in Ballard's E - i plot is 0.08 ohm cm 2. The performance of relatively high power

densities (IA/cm 2 at 0.7V) in solid polymer electrolyte fuel cells with low platinum loading

1 Exchange Graduate Student from Ruhr University, Bochum, Germany, FR.

95
03

I- 0

LLI

80-

LU

0

,=,,

<

,0i,

"-I

_. -._ o

(1) 0 _ _.."

-o _ o r- (Do o). 0 , _0.c: _-.

__ _.-- ---_ _

_o_=o °;{ 0 "-'

>,_ =

_R8_o

_-m _: _'_

o) l.U _

_E

= -_o

t-- o¢.-

_o o_

_ o_

q0 co

'- "" a:= 1=

o-o :_ o

_Eo_ E

_<o._ =

•,-_ ¢6

g6

l

1
electrodes was first demonstrated at Los Alamos National Laboratory by Srinivasan, Ticianelli,

Derouin and Redondo (5-7).

The solid polymer electrolyte fuel cell system has major advantages over the alkaline fuel

cell one in terms of lower operating temperature, ability to start from cold, and total lack of

electrolyte management problems. From a cold (10°C) start-up, this system can be designed to

give full power in under 30 seconds. The fuel cell system can easily be programmed for instant

start-up by arranging a stack hydrogen dead-volume after the cut-off valve equal to twice the

corresponding oxygen cut-off-dead-volume. Short-circuiting will rapidly reduce hydrogen and

oxygen partial pressures to close to zero, with production of water that is sealed into the stack, and

so will maintain the membrane in peak condition. Thus, a cold solid polymer electrolyte fuel cell

stack can be safely and indefinitely left on open circuit, if prevented from freezing, and it can still

be ready for rapid start-up. This is a great advantage compared with the high-performance alkaline

fuel cell system, which must be maintained at the required minimum temperature to prevent

electrolyte solidification, and whose open-circuit storage under these conditions is doubtful.

1.2 N.eccssary Electroc!.e Kinetic Criteria for Attainment of High Power Densities

The heart of the fuel cell system is the electrochemical cell. Even though the electrode

kinetics of the electrochemical cell have been adequately dealt with in books (8,9) and review

articles in journals (10), it is worth recapitulating the electrode kinetic criteria for attainment of high

power densities. As will be seen from the brief analysis presented below, the electrode kinetics of

fuel cell reactions pose the critical issues and problems.

Firstly, there is no doubt that the only fuel ceils which are capable of attaining high power

densities are the ones using hydrogen and oxygen as reactants. Secondly, it has been clearly

demonstrated that only the fuel cells with alkaline and perfluorinated sulfonic acid polymer

electrolytes are capable of attaining high power densities. The main reason for this is that there is

hardly any anion adsorption on electrocatalysts from these electrolytes and thus the poisoning of

the oxygen reduction reaction is minimal. Thirdly, in fuel cells with both these electrolytes, mass

transport limitations are not visible at current densities up to several A/cm 2. Under these

conditions, the cell potential (E) - current (i) relations may be expressed by the equation:

E =Eo- blogi- Ri (1)

In perfluorinated sulfonic acids and their polymeric acids (for e.g., Nation), the kinetics of

hydrogen oxidation is extremely fast and the potential of the hydrogen electrode varies linearly with

current density up to high current densities (a few A/cm 2) (7). In alkaline electrolytes, this is not

the case because the exchange current density for this reaction is considerably less than in the

perfluorinated sulfonic acids(11). However, the oxygen reduction reaction is faster in alkaline

medium than in perfluorinated sulfonic acids. In both media, the relation between the oxygen

electrode potential and current density is semi-logarithmic, and this accounts for the second term in

the right hand side of Eq. 1. In this equation, b represents the Tafel slope for the oxygen reduction

reaction. The first term on the right hand side, may be further expressed by:

Eo = Er- b log io (2)

where Er is the reversible potential for the cell and io is the exchange current density for the oxygen

reduction reaction. Differentiating equation (1), one obtains:

dE b

--= - .-- - R (3) di 1

97
At low currentdensities,the first term on the right handside of Eq. 3 is predominant and is

reflectedin the high slope(which graduallydecreases)of thecell potential- currentdensityplot

(seeFig. 1). At higher current densities, the second term becomes important and is responsible for

the linear region in this plot. Contributions to R are from (i) the ionic resistance of the electrolyte;

(ii) electronic resistance of the electrodes; (iii) charge transfer resistance of the hydrogen electrode;

(iv) charge transfer resistance of the oxygen electrode which is due to the difference in the

activation overpotential between the two current densities over which the slope was measured; and

(v) any small contributions to mass transport resistances. As seen from figure 1, if mass transport

becomes important there will be a rapid fall-off of cell potential with current density and thus only

close to the limiting current density does mass transport control exert its influence. This

apparently simplified equation applies over 3 decades of current density ( lm A/cm :_to 2 A/cm z)

for high power density solid polymer electrolyte fuel cells (see Section 3).

It is worthwhile at this stage to focus on the electrode kinetic requirements to attain high

power densities by an examination of Equations (1 to 3) and Figure 1. Two factors clearly

dominate the shape of the cell potential - current density relation: these are the Tafel parameters for

the oxygen reduction reaction and the ohmic overpotential in the cell. The highest values of the

exchange current density for oxygen reduction in porous gas diffusion electrodes is approximately

10 -6 A/cm 2. The minimum value of R in Eq. 1 which has been reported is 0.05 ohm cm 2 for the

alkaline fuel cell system (operating at 150°C and 8 atm) and 0.08 ohm cm z for the solid polymer

electrolyte fuel cell system (operating at 95°C and 5 atm). In alkaline electrolytes, the Tafel slope

for oxygen reduction on the best electrocatalyst (90% Au 10% Pt) is 0.04 V/decade, whereas in

solid polymer electolytes it is 0.06 V/decade. The best reported cell performance in the alkaline

fuel cell system is at a temperature of 150°C and 8-10 atm pressure. In the solid polymer

electrolyte fuel cell system, most of the work has been carried out at close to 100°C, whereas the

best, reported, performance by Ballard Technologies, Inc. is at 120°C. Use of a value of Er = 1.20

V is a close approximation for the reversible potential of the fuel cell in both these environmentS.

Thus at a current density of 3 A/cm z, and a cell potential of 0.8 V(an expected performance of fuel

cells for defense applications), one may write the following expressions for the exchange current

densities of the oxygen reduction reaction:

Alkaline fuel cell:

0.8

SPE_fuel cell:

0.8

= 1.20 + 0.04 log io - 0.04 log 3 - 0.05 x 3

•". io = 1.7 x 10 -6 A/cm 2

(4)

= 1.20 + 0.06 log io - 0.06 log 3 - 0.10 x 3

.'. io = 6.5 x 10 -2 A/cm 2

(5)

This simplifiedarialysis signifies that the exchange current density for oxygen reduction in alkaline

electrolyte must be equal to or greater than about 2 x 10 -6 A/cm z (based on the geometric area of

the electrode) to reach the specified goals. On the other hand, with the value of R (0.1 ohm cm 2)

assumed for the solid polymer electrolyte, the exchange current density for the oxygen reduction

reaction will have to be greater than 7 x 10 -2 A/cm 2 (again based on the geometric area of the

electrode). If, however, we were to assume the same value of R for the solid polymer electrolyte

fuel cell as that for the alkaline fuel cell, the exchange currenl density for the oxygen reduction will

have to be 2 x 10 "4 A/cmL Considering the exchange current densities, which have been reposed

for oxygen reduction on smooth surfaces in alkaline (io = 10 .8 A/cm 2) and fluorinated sulfonic acid

(io = 10-7 A/cm 2) electrolytes at room temperature and assuming (i) an increase of these values by a

factor of 10 for operation at elevated temperatures (150°C for the alkaline fuel cell system and

7

98
120°Cfor theSPE®fuel cell system);and(ii) aconservativevalueof 100 for the roughness factor of

the porous electrodes, it should be possible to reach the desired goal of operation at 3 A/cm 2 and

0.8 V/cell, only if the slope of the linear region of the cell potential - current density is 0.05 ohm

cm 2 or less.

1.3 A Snap-Shot Version of the Status of Solid Polymer Electrolyte Fuel Cell Technology

In three recent articles (5-7) the present status of the solid polymer electrolyte fuel cell

technology is summarized. For detailed information, the reader is referred to these publications.

In this paper a "snap-shot" version (Table 2) of the progress made in this field is presented.

Table 2 demonstrates that significant progress has been made during the 1980's in solid polymer

electrolyte fuel cell technology in spite of the fact that the funding for this program is miniscule

compared to the investments made for the development of alkaline, phosphoric, molten carbonate

and solid oxide fuel cells. In hardware production of solid polymer electrolyte fuel ceils, Ballard

Technologies Corporation and Ergenics Power Systems, Inc. have done extremely well. Both are

at a stage where they can custom-make 1 1/2 kW power plants. However, these systems, as well

as the previously developed ones by General Electric Company, contain electrodes with high

platinum loadings (4 mg/cm2). The first demonstration of high power density solid polymer

electrolyte fuel cells with low platinum loading electrodes was by one of the authors in this paper

(SS) and his coworkers at Los Alamos National Laboratory (LANL); a brief description of this

work is presented in the next sub section.

1.4 A Synopsis of Advances in the Attainment of High Power Densities in Solid Polymer

Electrolyte F_el Cells with LOW Platinum Loading Electrodes

Since the first Space Electrochemical Research and Technology (SERT) Conference, in

which the results of the first set of investigations on solid polymer electrolyte fuel cells with low

platinum loading electrodes were presented(5), there has been considerable progress in developing

methods for the attainment of high power densities (about 0.7 W/cm 2) in such types of fuel cells.

It must be noted that the Nation (registered trademark of Dupont) membrane was used as the

proton-conducting membrane in all these investigations. The methods used to attain the high

power densities may be summarized as follows:

(i) optimization of the amount of Nation impregnated into the electrode structure;

(ii) hot-pressing of the Nation-impregnated electrodes to prepurified Nation membranes at

120°C (close to glass transition temperature) and 50 atm pressure;

(iii) optimization of the humidification of the reactant gases at a temperature of 5°C for

oxygen or air and 10-15°C for hydrogen above the cell temperature;

(iv) operation of the cell at elevated temperatures and pressures (say, 80 ° C and 5 atm); and

(v) localization of platinum by fabrication of electrodes with a higher percentage of

platinum crystallites on high surface area carbon (i.e., supported electrocatalysts with

20 wt% Pt/C rather than 10%) while still maintaining the amount of Pt in the electrode

(0.4 mg/cm 2) and by sputter-deposition of a thin f'dm of platinum on the front surface

(0.05 mg/cm2), corresponding to 500/_ film a smooth surface).

By use of all these methods, it was possible to attain a current density of 1 A/cm 2 at a cell

potential of 0.64 V with H2/O2 as reactants and 0.580 V with H2/air reactants at 80°C and 3/5 atm

pressure (3 on hydrogen and 5 on oxygen side).

The cell potential (E) - current density (i) plot fitted the equation (1). The electrode kinetic

parameters for the cell Eo, b,and R were calculated using a non-linear least squares fit of this

equation to the experimental points. The slope of the Tafel line for oxygen reduction was found to

be 0.050 to 0.060 V/decade and was independent of temperature.

99
0

-1

_Z

0 o

_l,.a

_,, [.ra

o

_o_

0

•_.=.o= _ _

_'_

.-.I g

"< "ag

0

c

-___

.o _._

M _'_

_o e_

0

o

"g::l

_ ..=-=

Ii_ i... ,fat

o 0.,., _ 0

_.._'_"

E.,_ _._ _

B',::J o_ ..o

'._" _ :-.inl _ I= _

"_ o _...o a

,-,_ "_ _ _,'_

o _;_ __,o_'_

O_.a • _ _._

,-. R o _a = _al=

¢.1

_g

,_

,_I_ ,

_

_>

•-, _ g:l 0

£_--=

g

o>o =I

_o,..

b_O

e4

88_

mo.o

ml=

_'_ _

C_

4 ,6

z_

t,_ v

o_

_o

,,4

lOO
Since small variations in the parameter b give rise to misleading values of io, it was decided

to use the current density at a cell potential of 0.900 mV (at this potential mass transport and ohmic

effects are negligible and so also is the overpotential at the hydrogen electrode) as a measure of the

electrocatalytic activity of the electrode for the oxygen reduction reaction. It is extremely important

to have this value as high as possible to minimize the steep decrease of slope of the cell potential -

current density region at very low values of curent densities. In the best case, i.e., with H2/O2 as

reactants at 3/5 atm and 80°C in cells with electrodes composed of 20% Pt/C supported

electrocatalysts and 0.4 mg Pt/cm 2 on to which a thin film of Pt was sputtered, the current density

at 0.900 V was 160 mA/cm 2. Further, the calculated value of R (0.24 ohm cm 2) was quite close to

the high frequency resistance (0.21 ohm cm2), indicating the absence of mass transport

overpotential over the measured current density range. The increase of cell potential with pressure,

at constant current density, showed the expected semi-logarithimic dependence. The cyclic

voltammetric technique was used to determine the electrochemically active surface area of the

electrode. By use of this method, the platinum utilization of the electrode with the best

configuration of electrodes (20% Pt/C 0.40 mg Pt/cm2 plus sputtered film (0.05 mg Pt/cm 2) was

found to be about 15-20%. It was concluded in the second publication (7) since the SERT

conference that to increase the power densities further, it is necessary to use membranes with

higher specific conductivities and better water retention characteristics.

1.5 Scope of Present Work

The work described in the preceding sub-section was carted out by one of the authors (SS)

of this paper and his former colleagues at LANL. This work has since then been continued in our

laboratories with the objectives of increasing the power densities still further, as required for some

applications (for example, silent mobile power source for transportation). While the work at

LANL was focussed on developing on solid polymer electrolyte fuel cells using reformed

hydrogen (steam-reforming of methanol), the work at our Center (CESHR) is concentrating on

developing fuel ceils using pure H2 and 02. As seen from the results in the preceding section, the

"high priority" tasks which are envisioned to beimportant are (i) to develop methods, alternative to

sputter-deposition to localize platinum near the front surface of the electrode and (ii) to reduce the

ohmic overpotential in the cell. The reason for importance of the former is that sputter-deposition

is not economically feasible compared to wet-chemical methods, which can be automated, for the

large-scale modification of surfaces of electrodes for fuel cell systems with a reasonable power

output (20 kw and higher). The methods, which are examined in this work for the deposition of a

thin layer of Pt on the front surface of the electrodes are: (i) brushing the surface with a

chloroplatinic acid solution (desired quantity and drying); (ii) brushing surfiice with platinum black

particles suspended in Nation solution; and (iii) electrodeposition of Pt. Two approaches were

taken to reduce the ohmic overpotential in the cell: use of (i) thinner membranes and (ii) Dow

membranes with better conductivity and water retention characteristics.

2 EXPERIME.. NTAL

2.1 Wet Chemical/Electrochemical Methodz to Localize Platinum near Front Surfaces

The membrane and electrode (M&E) assemblies were fabricated using Protech porous gas

diffusion electrodes (20% Pt/C, 0.4 mg/cm 2) and proton conducting polymer membranes (Dupont

and Dow Chemical). Localization of an extra layer of platinum on the active surface of these

eleclrodes was achieved by wet chemical methods: (i) brushing with Nation solution containing

unsupported Pt catalyst; (ii) chemical deposition of Pt from H2PtC16 and (iii) electrochemical

deposition of Pt also from H2PtCI6.

101
In the first method, 0.5 mg of fuel cell grade unsupported Pt catalyst (Prototech) was

sonicated in 5% Nation solution (0.6 ml) and two electrodes (each of 5 cm 2 area) were brushed

with the mixture, using a soft printing brush, as uniformly as possible. The brushed electrodes

were next dried in a vacuum oven at 70°C and then impregnated with Nation in the manner

described in previous publications (5-7), prior to fabrication of the M&E assembly. Chemical

deposition of Pt (0.5 mg/cm2) was carded out by brushing the active surface of the electrodes with

an appropriate amount of 1% H2PtCI6 solution in 0.05 M H2SOn,-_-Ifited-with an equaI'Voiumeof

methanol; this was then followed by heat treatment of the electrodes fo( tqaree hours 'at 200_n

air. Electrodeposition of Pt, on the active side of Protech_electrodes using the same H2PtCI6 (I %)

- methanol (1:1 vol.), was carded out at a constant current density of 2 mA/cm 2 to deposit -0.05

mg/cm 2 of Pt; a gold foil was used as anode during electrolysis.

2.2 Assembly of Single Cells and Performance Evaloation

..... The main components ofa_single cell assembly _e:as described in previous papers(5-7).

The essential components are two graphite plates, the M&E assembly and gaskets containing the

gas inlet and outlet and ribbed channels for the distribution of reactant gases behind the porous g_ts

diffusion electrodes. The M&E assembly is positioned between two such graphite piates. The

reference electrode (a small piece of fuel cell electrode) is embeded in the graphite plate on the

anode side. Teflon coated fiberglass cloth gaskets are placed between the membrane and each of

the graphite current collectors to prevent gas leakage and to avoid excessive compression of the

electrodes. The current collection from the fuel cell uses a copper plate positioned behind each of

the graphite plates. Compression of the overall assembly is made by means of two 1/4 inch thick

stainless steel end plates and four bolts. Teflon sheets are used forelectric insualtion of the single

cell and the end plates. Though incorporation of a single cell in the test station with all its

peripheral components (e.g., teffiperature Crfitrollei', humidification Chamber, fl0W meters, back

pressure regulator) have also been described in the referen-ced p-fibliCafiofi-S :, theE0rnp0nents ofthe

single cell and the test station are illustrated in Figs. 2&3 respectively. The performance evaluation

was carried out using a microcomputer (IBM - XT) a power supply (HP 6033A) and a data

acquisition unit (I-IP-3421A) interconnected through a GPIB bus. During performance evaluation,

the fuel cell is connected in series with the power supply for its operation under galvanostatic load

and cell potentials as well as half cell potentials versus current density data Were recorded

periodically. Measurements were made at different temperatures and pressures of reactant gases

(H 2 at the anode and air/O2 at the cathode). _

3 RESULTS AND DISCUSSION

31J AComparis0n of the Effects of the S@tter-Depos_cin-_ind _ Wet-Chemical/'

Electrochemical Methods on the Performance of Single Cells

As stated in Section 1, one of the main objectives of this work was to find alternatives to

sputter-deposition to localize the platinum near the front surface of porous gas diffusion electrodes

with low platinum loading electrodes. The cell potential current density plots (Fig. 4) show that

fuel cells with the sputter-deposited and chloroplatinic acid treated electrodes exhibit similar

performances throughout the entire current density range (1 to 2000 mA/cm2). This is not so in the

case of the of the fuel cell with electrodes on to which a thin layer of Pt black was deposited. A

close examination of the linear E vs i (Fig. 4), as well as of the semi-logarithmic (E+ i R) vs log i

(Fig. 5) plots shows that the oxygen electrode is only activation-controlled up to a current density

or 1 A/cm 2. Fig. 5 and Table 3 illustrate that the fuel celI with H2PtC16 treated electrodes sh0wsa

higher oxygen electrode performance than the electrodes wlt_t_e Pt=sputter-deposited and

Pt/Nafion thin layers. This reflects a higher surface area of the H2PtCl6 treated electrode. The fuel

cell with the Pt black treated electrodes have an unusually high slope in the linear region of the E-i

m

z

h

102
TABLE 3

Effect of Method of Deposition of Thin Layer of Platinum (0.05 mg/cm 2)

on Front Surface of Prototech Electrodes (20% Pt/C, 0.04 mg Pt/cm 2)

on Electrode Kinetic Parameters (see Eq. 1) for H 2/02 Solid Polymer

Electrolyte Fuel Cell with 100 I_m thick Nafion Membrane.

Electrode Kinetic Current

Parameters for Cell Density

Membrane Type Cell Pressure Eo b R at 0.9V

and Thickness Temperature H2 /O2 V V £_J'cm2 mA/cm 2

Sputtered Pt 50 1/1 0.930 0.044 0.339 7

" 50 4/5 0.993 0.047 0.301 35

" 70 4/5 0.983 0.040 0.252 46

" 85 4/5 0.999 0.450 0.201 52

" 95 4/5 0.997 0.500 0.189 61

Unsupported

Pt/Nafion, Brushed 5 0 1 / 1 0.964 0.052 0.501 1 2

50 4/5 0.956 0.048 0.452 48

" 70 1/1 1.024 0.059 0.505 10

" 70 4/5 1.014 0.052 0.462 51

" 85 4/5 1.008 0.045 0.414 61

" 95 4/5 1.005 0.044 0.368 35

H 2 PtCI 6 in

H20-CH3OH-br ushed 50 1/1 0.978 0.056 0.338 19

50 4/5 1.024 0.053 0.341 72

85 415 1.019 0.044 0.253 80

95 415 1.020 0.056 0.189 83

103
plot and the departure from lineadty occures at a relatively low current-density. The slopes of the

linear regions of the E-i plots in the other two cases are similar for identical operating conditions

but not low enough to reach the desired goals of power density (See Section 1.2). However, it is

encouraging to note that the wet chemical method of application of chloroplatinic acid followed by

the heat-treatment provides electrodes which exhibit similar performances to the sputter-deposited

ones.

3.2 Effects of Thickness of Membrane on Performances of Fuel Cells

Most of the LANL studies reported in references 5 to 7 used Nation 117 membranes. Fuel

cells with these membranes yield a slope of approximately 0.25 ohm cm 2 in the linear region of

the E-i plot. If the resistance of the test cell fixtures (0.05 ohm cm 2) is subtracted from this value,

the slope of the E-i (Table 3) plot (linear region) will be reduced to 0.20 ohm cm 2- This value is

still too high to reach current density of 2 to 5 A/cm 2 at reasonable cell potentials (> 0.._V)I Thus,

in this work experiments were carried out with Nation membranes having a thickness of 50

andl00 t.tm and the results of fuel cell performance compared with those in fuel cells with the

Nation 117 membrane (thickness: 175 gm). The 100 ktm membranes were of the "sweded type"

(sweding was carried out to roughen the surfaces of the membrane and make it of uniform

thickness). The performances of the fuel cells with the 50, 100 and 175 I.tm membranes are shown

in Fig. 6. While the fuel cell with the Nation 117 membrane begins to show mass transport

limitations at a current density of 1 A/cm 2, this is not the case for the single cells with thinner

Nation membranes. The single cell with 100 i.tm membrane shows mass transport limitation at 1.8

A/cm 2 while the cell with 50 gm membrane shows no mass transport limitation up to the highest

current density (2 A/cm2) of measurements.The electrode kinetic parameters extracted from these

experimental results and using equation (1), are shown in Table 4. The agreement in the "pseudo i

R" (all forms of overpotential which show a linear variation of potential with current density -

mostly ohmic ) corrected Tafel plots Fig. 7 for oxygen reduction in the cells with the 100 and 175

I.tm thick membranes are as expected. The sweded membrane (100 gm thick) shows a slightly

higher resistance than expected on the basis of its thickness as compared to the thicker (175 i.tm)

membrane (Table 4).

3.3 Dow Membranes: The Solution to Attainment of Super High Power Densities

In the work to date the highest power densities were attained by using Dow membranes as

the electrolyte layer. A comparison of the performances of single cells using the Nation and Dow

membranes is illustrated in Fig. 8. The Nation Membranes had thicknesses of 175 and i00 I.tm

while the Dow membranes were 125 I.tm thick. It is natural to expect a thinner membrane to have a

lower resistance but even if we take this into consideration the Dow membrane has a lower specific

resistivity (i.e., higher specific conductivity) than Nation. The reason for the higher specific

conductivity of the Dow membrane than that of Nation is that the monomer of the former may be

represented by:

while that of the later is:

CF2= CF O CF2 CF2 SO3 H

CF2= CF O CF2 CF O CF2 CF2 SO3 H

I

CF3

(6)

(7)

Hence, there are a greater number of sufonic acid groups in Dow polymer membrane than in the

Nation polymer membrane. Expressed in another manner, the Dow proton conducting polymer has

a lower equivalent weight than Nation.

104

-7

I
TABLE 4

Effect of Thickness of Nation Membrane and of Type of Membrane on

Electrode Kinetic Parameters of H 2 /02 Solid Polymer Electrolyte

Fuel Cell with Prototech Electrodes (20% Pt/C, 0.40 mg Pt/cm 2)

on to which a Thin Layer of Pt (0.05 mg/cm 2 ) was sputtered.

Electrode Kinetic

Parameters for Cell

Membrane Type Cell Pressure Eo b R 2

and Thickness Temperature H2 /02 V V Q/cm

Current

Density

at 0.9V

mA/cm 2

Nation 117,

1751J.m 50 1/1 .946 .049 .475

50 3/4 1.003 0.048 .428

80 3/4 .951 0.037 .437

85 3/4 1.007 0.046 .298

7

50

61

70

Sweded Nation

117, 100p.m

if

w

50 1/1 0.964 0.052 0.339

50 415 0.956 0.048 0.301

70 4/5 1.024 0.059 0.252

85 4/5 1.008 0.045 0.201

95 4/5 1.005 0.044 0.189

12

48

10

61

35

Nation 117,

50p.m

Dow Membranes

N

w

50 111 .833 .061 .188

50 1/1 .955 .061 .195

50 4/5 .995 .053 .157

70 1/1 .912 .062 .153

70 4/5 .994 .053 .116

85 4/5 1.002 .053 .111

95 4/5 1.000 .549 .110

<1

7

42

1.6

43

62

54

105
Figure 8 shows that the cell-potential current density plots overlap in the low current

density region (say up to 50 mV). This means that the activation-controlled behavior is unaffected

by the characteristics of the membrane (Dow vs Nation, Nation thick vs thin). One can easily

interpret this result on the basis that the electrochemical reaction occurs to a large extent within the

pores and the surface of the electrodes which are coated with the proton-conducting membranes.

With increasing current density, the role of the membrane (type, thickness) is dominant and

consequently the slopes of the linear region are different. One significant result is that the slope of

the linear region in the cell with the Dow membrane (thickness 175 _m) is less than that in the cell

with the Nation membrane (thickness 100 l.tm) even thaough the thickness of the Dow membrane

is greater than that of the Nation membrane. This result can be explained by the fact that the Dow

membrane has a higher specific conductivity than that of the Nation membrane.

There is a second interesting observation in the E-i plots. The departure of the E-i plot

from linearity at the higher current density appears to depend on the thickness of the membrane (cf

the results with Nafion membrane 175 and 100 I.tm thick membrane) and the type (cf Dow

membrane 125 ktm vs Nafion 100 i.tm). Departure from linearity at the high current densities are

associated with mass transport controlled processes. Generally, these are due to mass transport

limitations of reactants reaching the active sites in the electrode or of products (or inert gases like

N2 when air instead of oxygen is used) away from the electrocatalytic sites. However, the

electrode structures and conditions of humidification of reactant gases were identical in the three

cases. It is very likely that transport processes within the membrane can be rate-limiting. The

transporting species are protons, which migrate from the anode to the cathode under the influence

of the electric tield, and water molecules which are Carded with the protons. It is estimated that 3-6

water molecules are transported with the protons. Due to the resulting concentration gradient of

water molecules in the cell during operation one can expect water molecules to diffuse from the

cathode to the anode. Mass transport limitations could occur due to any one or more of these

processes. The Dow membrane shows a linear behavior up to a current density of 2 A/cm 2. This

is not the case in the cell with the Nation membrane which is slightly thicker than the Dow

membrane. This result lends insight on different types of proton conduction (say Grotthus type in

a Dow membrane vs classical proton transport through the electrolyte in Nafion). It is interesting

to point out that in previous studies it was noted that when mass transport limitation begins on one

electrode, it also sets in on the other at the same current density. A detailed modeling analysis of

the mass transport processes is necessary to interpret the result in the higher current density range.

The electrode kinetic parameters for the cells were calculated as described briefly in section

1.2. The parameters E0 and b are unaffected by the membrane. However, R, which represents the

rate of increase of cell potential in a linear manner with the current density, depends on the

membrane type and its thickness. The most significant contribution to R is the ionic resistance of

the membrane, as has been shown previously using high frequency measurements (7). The

striking result is that R is a factor of two less for the cell with the Dow membrane (thickness 125

_m) than for the one with the Nation membrane. This result follows from the fact that the Dow

membrane has a lower specific resistance. The calculated value of R is always slightly less than

the slope of the linear region in the E vs i plot. For example for the ceil with the Dow membrane

operating at 95°C, the R is 0.110 while the slope of the linear region is 0.113. The reason for this

is that the apparently linear slope of the E-i plot includes a small contribution of the charge transfer

resistance of the oxygen reduction reaction. In the cells used in this work the test cell fixtures have

a resistance of 0.05 ohm cm 2. If this is excluded from the values of R and of the slope of the E-i,

line the resulting values are only 0.06 and 0.08 ohm cm 2 respectively. The latter value is the same

as that of the slope of the E-i plot in the best Ballard Cells. Using the calculated values of R, the (E

+ Ri) vs log i plots (Fig.9) were made for the cells with the Dow (125 ktm) and Nation (100 I.tm)

membranes operating at 95°C and 4/5 atm. The Tafel behavior is observed for the former cell over

the entire current density range, whereas for the latter, there is departure from the Tafel line at the

same current density as the one in the E-i plot where it departs from linearity.

106
Another resutl is worth mentioning. The half and single cell potentials are plotted as a

function of current density (Fig. 10) for the cell with the Dow membrane. The hydrogen electrode

exhibits a linear behavior throughout the entire current density range. The hydrogen overpotential

at 2A/cm 2 is only 20 mV. Thus the charge transfer resistance due to the hydrogen electrode is only

0.01 ohm cm 2 which is a contribution of this electrode to the calculated R value or the measured

slope of the cell potential vs current density plot.

4 CONCLUSIONS

The conclusions which may be drawn from these studies are:

(i) The chloroplatinic acid method of treatment of the electrodes provides a satisfactory

alternative and considerably more economic method than sputtering for the deposition

of a thin layer of Pt on the front surface of the electrodes, which is essential for the

attainment of high power densities.

(ii) Use of thinner membranes is advantageous from the points of view of lowering the

ionic resistance and lowering mass transport limitations at higher current densities.

(iii) Striking results have been obtained in cells with the Dow membrane which has a

higher ionic conductivity, lesser mass transport limitations (H + and H20) and better

water management characteristics than Nation.

(iv) The performances of a solid polymer electrolyte fuel cells with low platinum loading

electrodes are approaching those of the ones with ten times the platinum loading in

respect to attainment of high power densities.

ACKNOWLEDGEMENTS

This work was supported by a contract from DARPA/ONR, Code 1113 ES. The authors

are grateful to Dow Chemical (via Drs. Glenn Eisman and Jeffrey Gunshor), Freeport, Texas, for

providing Dow polymer membranes used in this work.

REFERENCES

1. (a)

(b)

J. K. Stedman, 1985 Fuel Cell Seminar Abstracts. P. 138. Courtesy Associates,

Washington D. C. (1985).

R. Warnock, Wright-Patterson Air Force Base, Personal communication.

. (a) D. Watkins, D. Kircks, D. Epp and A. Harkness, Proceedings of the 32nd Annual

Power Sources Conferences, Cherry Hill, N.J., June 9-12, 1986. The Electrochemical

Society, Inc. 590, (1986).

(b) D. Watkins, Ballard Technologies Inc., personal communication, (1987).

° O.J. Adlhart and M.J. Rosso, Ergenic Power Systems, Inc.,Wyckoff, NJ, Personal

communication.

, K. Strasser, Siemens, Erlangen, Germany FR, Personal communication.

. S. Srinivasan, E.A. Ticianelli, C.R. Derouin and A. Redondo, J. Power sources,

22. 359 (1988).

107
,

°

.

o

E. Ticianelli, C.R. Derouin, A. Redondo and S. Srinivasan, J. Electrochem. Soc.,

135, 2209 (1988).

Edson, A. Ticianelli, Charles R. Derouin and Supramaniam Srinivasan, J.

Electroanal Chem., 25___!,275-295 (1988).

J. O'M Bockris and S. Srinivasan, Fuel Cells; Their Electrochemistry, McGraw

Hill Publishing Company, New York (1969).

A. J. Appleby and F. R. Fowkes, Fuel Cell Handbook, Van Nostrand Reinhold,

New York (1988).

S. Srinivasan, J. Electrochem. Soc., 136.41C (1989).

M. A. Enayatullah and J. O'M. Bockris, Proceedings of the Symposium on

Electrode Materials _and Processes for Energy Conversion and Stbrage, S.

Srinivasan, S. Wagner and H. Wroblowa (eds.), Proceedings Volume 87-12 256,

The Electrochemical Society, Pennington, NJ (1987).

108
1.5 1.0

Figure 1.

,_ 1.0

4J

0.5

a

THF_RMODYNAMIC

REVERSIBLE CELL

POTFJCTIAL (E)

THE IDEAL CELL POTEN_AL-CURRENT RELATION

CELL POTENTIAL LOSSES DUE TO

\_"" ACTIV IATION OVERPOTENT[AL

I _(LACK OF ELECTROCATALYSIS) LINEAR DROP IN CELL

/ POTENTIAL MAINLY DUE TO

__._ _ OHMIC LOSSES IN SOLUTION

MASS TRANSPORT

LOSSES CAUSE OF'

DECREASE OF CELL

POTENTIAL TO ZERO

I !

0 0.5 1.0 0

CELL _, arrlDeres

INTRINSIC

MAXIMUM

_ EFFICIENCY

0.5

Typical plot of cell potential versus current for fuel cells illustrating regions of

control by various types of overpotentials.

Copper Plate

Current

/Collerlors_

Anode Cathode

S S Endplete

;ket ,/

ELECTROCHEMICAL REACTIONS

SPE MEMBRANE

Figure 2. Schematic of solid polymer electrolyte fuel cell assembly.

]09
"o i _ "

i E ! _IN

T

• -'t"

.,Hi

I Compuler J

I

I

I Data

Acquisition I

Temperature, c_nt_o(|ers

r ......... %°°.°°

s o_' g

b

,.f,

tEL | CE.L_.

I

-o

o

L__.

Figure 3. Schematic of solid polymer electrolyte single cell test station.

1.2

1.1

I

O.9

0.11

0.5

0.4

0.3

0.2

0.1

0

Figure 4.

I I I I I I t I I I I I I • i I ! 1 ! I

0 200 400 ISO0 800 1000 1200 1400 1600 1800 2000

Cell potential vs. current density plots for single cells with electrodes having

localized Pt layer deposited by different methods, all operating at 95°C and

with H2/O 2 at 4/5 atm. Sputter deposited (-), electrochemically deposited

(+),chemically deposited from H2PtCI 6 (A) and brushed with Nation solution

containing unsupported Pt i¢,). Pt loading on each electrode: 0.45 mg/cm2.

110
w

÷

(2.

8

Figure 5.

1.2

1.1

1

O.g

_ 0.7

0.6 0.5

_ 0.4

O.3

O.2

0.1

0

Figure 6.

,I.2.

H2/02 4/5 Af.m Ce]] Temp. S;5*C

1.1-

1.0- .90.

.80.

o

.7_.

o

.60-

i

• 50-

.30_

+ - H,?.PtCI6 Drushed

• 20

O - unsupported Pt: brushe¢/

.10

N - Pt: sputtered

J.O 100 t000

Current: Density [=A/cm'2]

Plots of (E + JR) vs log i for single cells with electrodes having localized Pt

layer deposited by different methods, all operating at g5°C with H2/O 2 at 4/5

arm. sputter-deposited (_), chemically deposited from H2PtCI 6 (+), and

brushed with Nation solution containing unsupported Pt (o). Pt loading on each

electrode: 0.45 mg/cm2.

ORIGINAL PAGE IS

OF POOR QUALITY

0 2oo 4oo 6oo eoo i_ 1200 1_ _600 1_

cu_ ee_s_ (m,V'='n"2)

0 IO0 + 50 _ 175)Jm

Cell potential vs current density plots for single cells with Nation 117

membranes of two different thicknesses, operating at 95°C with H2/O 2 at 4/5

eachatm" electrode:Nati°n' thickness:1750.45 mg/cm_m. (_,100 p.m (=) and 50 p.m (+) Pt loading on

111
l.l'

.S0'

.80.

.70-

>

.00

e.

.i,

.50

.40+

#.

.., .30,

.20.

.10"

H#./02 4/5 At.m Ce|| Tamp. 05"C Nsflon Nelm.

0 - 0.t75 as Hemprane

, - 0.100 M Xembrmne

t0 4100

: : : :::':

4000

L_Irrlnt Density [i£/¢m'2]

: : " ::::"

Figure 7. Plots of (E + iR) vs log ifor single cells with Nation membranes of two different

thicknesses, operating at 95°0 with H2/O 2 at 4/5 atm. Nation117, thickness:

175 I_m (o) and 100 llm (-_). Pt loading on each electrode: 0.45 mg/cm 2.

Figure 8.

121.1 _ H2/02, 4/5 arm

1

o.g

>_ 0.8

nu 0.7

0.6

0.5

0,4

0.3

O,2

0.1

0

0 200 400 600 800 1000 1200 1400 11500 1800 2000

_m,te,n"De,tarry(mM,:m"_)

a Bow 125/Rn -I- NAFION 100/_n <_ NAFION 175 k_n

Cell potential vs current density plots for single cells with different membrane

materials, operating at 95°C with H2/O 2 at 4/5 atm. Dow membrane, 125 p.m

(=), Nation 117, 175 p.m (G), and Nation 117 (sweded), 100 p.m (+). Pt

Ioding on each electrode: o. 45 mg/cm 2.

112
'1°_.

.!.1

1.0

.90

.BO

.70

_>

.BO- m

4"

.50.

.40-

,30-

.20"

,10.

DDW VS. NAF:_ HEMB_ANE sputtered Pt. Electrodes

H2/02 4/5 Ate Cell Temperature 95"C

0 - Dow Membrene 5 m[1

x - Nafton Hembpane 4 ell

`10 t0o to0o

Current Density [mX/cm'2]

Figure 9.

o

>

fJ

Plots of (E + JR) vs log i for single cells with different membrane materials,

operating at 95°C with H2/O 2 at 4/5 atm. Dow membrane, 125 pm

Nation 117, 175 _m (.)(-).Pt loading on each electrode: 0.45 mg/cm 2.

1.2

1.1

I

0.9

0.8

0.7

0.6

0.5

0.4

0,3

0.2

0.1

0

H 2 Electrode

0 200 400 600 BOO 1000 1200 1400 1600 1800

CURRENT DENSITY (mA/cm'2)

D CELL 95'C + 02 95'c ¢ H2 Io5"c

2000

(o),

Figure 10. Cell and half cell potentials vs current density plots for a single cell with Dow

membrane (thickness: 125 pm) and Pt-sputtered Prototech electrodes (PI

loading: 0.45 mg/cm 2) operating at 95 °C with H2/O 2 at 4/5 atm.

113

N90-20466

THE APPLICATION OF DOW CHEMICAL'S PERFLUORINATED MEMBRANES

IN PROTON-EXCHANGE MEMBRANE FUEL CELLS

G.A. Eisman

The Dow Chemical Company

Freeport, Texas 77541

Dow Chemical's research activities in fuel cells revolve around the development of

perfluorosulfonic acid membranes useful as the proton transport medium and separator.

The following work will outline some of the performance characteristics which are

typical for such membranes.

INTRODUCTION

The use of fuel cells as a power source has been successful for manned space applications on short-lived missions as is evident from the Gemini, Apollo and Orbiter missions. In the latter two programs, the alkaline technology has been the system of

choice; whereas the Gemini program experimented with the proton-exchange membrane

technology (PEM). Although the alkaline technology has enjoyed a certain degree of

popularity among the various fuel cell concepts for space applications, there are no

clear-cut favorites for one fuel cell technology over others when the duty cycle

becomes long as it would in extensive long-lived missions.

Until recently, the PEM fuel cell technology has been a laboratory curiosity except

for a few brief missions during the Gemini program. Although attractive for many

reasons, including long life (Ref. ]) and reduced failure rates as a result of pressure upsets (when compared to free electrolyte systems), the lower power and efficiency characteristics have kept the technology shelved. Recently, Dow Chemical

developed a new perfluorosulfonic acid proton transporting ionomer which can be

applied in fuel cells (Ref. 2-3). The results of tests with the Dow membrane as

reported in the literature (Ref. 4-7), have demonstrated that the materials have a

significantly reduced internal resistance. For fuel cells, this reduced resistance

translates directly to performance and efficiency since a membrane with this characteristic results in significantly enhanced proton transport rates with a lower emf

driving force.

The performance of a proton-exchange ionomer in a fuel cell is related to many factors, some of which include membrane ionic activity, hydration level, visco-elastic

properties, as well as electrode characteristics, cell design, and mode of operation.

This paper will be concerned with the properties of the Dow materials which are

related to performance characteristics of a PEM fuel cell system.

EXPERIMENTAL

The Dow experimental membrane utilized in this study is characterized by having an

ion-exchange capacity of I.]3 and being .013 cm in thickness All cell tests were

carried out in a single cell 50 cmc active area PEM fuel cell designed and built by

the Hamilton-Standard Corp. for Dow Chemical. The cell is a modified version of the

hardware delivered to Los Alamos National Laboratory under a transportation contract

with Hamilton-Standard (NASA Contract #P.O. No.g-x53-D6272-1).

The current interrupt experiments were carried out using a home-made fast response

switching device coupled to a Nicolet storage oscilloscope.

PRECEDING PAGE BLANK NOT FILMED

ll5
The membrane and electrode assemblies were fabricated using proprietary techniques.

The electrode loadings were 2 mg/cm_/side.

RESULTS AND DISCUSSION

Membrane Properties

The new perfluorosulfonic acid membranes recently introduced by Dow Chemical, as

referenced above, exhibit significantly lower IR losses resulting in membranes which

can be operated at ultra-high current (power) levels. This behavior is presented in

Figure I along with the state-of-the-art performance characteristics for a commercially available fuel cell membrane manufactured by DuPont under the tradename

Nafion®. As can be seen from the data, the Dow materials can generate substantially

higher current densities than systems utilizing the _uPont material. The voltage

losses which appear to be great at 3500 amps/ft _ (3771 amps/cm _) are actually a

result of using high resistant cell plates and hardware. It is approximated that the

loss due to the cell hardware is close to 200 mv.

Although themembrane is capable of higher current levels, the attractive characteristic and one which will be critical for long spac_ flights is the higher voltage at

the lower current densities, i.e., 100 - 500 amp/ft _. This enhanced voltage is a

direct result of the reduced internal resistance of the membrane which is attributed

mainly to the high sulfonic acid ionic content responsible for proton transport. The

advantages of such a higher voltage are many, but the most important is an increase

in the total electrochemical efficiency.

The Dow membrane is based on chemistry which leads to a reduced mass in the sidechain which is bonded to the long CF2 backbone. Presented in Figure 2 is the chemistry for the synthesis of the monomer and the copolymerization step with the tetrafluoroethylene; while in Figure 3, the resultant structure is depicted. Also presented in Figure 3 for comparative purposes is the structure of DuPont's Nafion.

The membrane properties pertinent for fuel cell operation are tabulated in Table ]

for the Dow fuel cell membranes. Such characteristics as ionic conductivity, water

permeability, gas permeation, and properties related to mechanical strength are

listed. The data presented in the table covers a range rather than single data

points because, at this time, Dow has prepared numerous membranes of different ion

exchange capacity which all have been successfully tested in fuel cell operation.

Fuel Cell Test Results

The hardware used during the membrane performance tests is a modified version of th

cell presented in Figure 4. The hardware utilizes carbon/Kynar plates with a 50 cm

active area. There is a heat exchanger on each side of the reactor compartment and

an internal humidification compartment.

Presented in Figure 5 is a polarization curve of the membrane previously described.

As can be seen from the figure, the membrane is capable of operating at high efficiencies at low current densities or lower efficienc_es at hig_ curreqt densities.

The resistance in the ohmic region is .220 ohm-cm ( (2.37X10 -_ ohm-ft'). This data

contains the resistance of the package and therefore is not an IR free number. It

should also be pointed out that the polarization data at each point was recorded only

after the system appeared to reach equilibrium at the given current density. Figure

6 is _ plot of the same data as Figure 5 but expanded in the region of 0 to 500

amps/ft _ .

116
Although steady-state performance can be achieved at any current density demonstrated

in the polarization data Rf Figure 5, the voltage of the cell vs. time at a single

current density (200 amps/ft _) for a relatively short period of operation is presented in Figure 7. These data include predetermined shut-down periods on a daily

basis. As can be seen from the plot, the voltage over the time element investigated

was invariant.

Current interrupting was carried out in order to determine the electronic resistance

effects of the hardware and remaining ionic resistance and over-voltage oK the membrane and electrode package. The tests were carried out at 100 amps/ft L (and less)

due to a limitation in the interrupting switching device. A typical result of the

tests is depicted in Figure 8 for the current membraneunder study. As can be seen

from the data presented in the figure, the IR free componentis approximately ]5 mv

at this current density and was measuredat approximately 50 microseconds after the

current break. This translates to an interqal resistance of 3 milliohms at 100

amps/ft_ (5 total amps) or .15 ohms-cm_. Whenthe IR free voltage (.933) at 100

amps/ft_ is subtracted from the open circuit potential (1.05 V) the resulting voltage

due to the membrane, membrane/electrode, and electrode overvoltages, amounts to

approximately 117 millivolts. In order to determine the package ohmic contribution,

the ce_l resistance is subtracted from the slope data. Such a calculation yields .07 ohm-cm (.222 - .15 ohm-cm_). A list of IR free values vs. at current densities

between25 and 100 amps/ft _ are presented in Table 2.

In order to determine the response times for the membraneand electrode assembly in

the Dowhardware utilizing a Dowperfluorosulfonic membrane,a series of tests was

carried out not unlike the above whereby the current was "interrupted" numerous

times. In these tests, the cell was allowed to reach equilibrium at the desired current density prior to the interruption. The cell voltage was monitored as a function

of time and recorded (Figure 9) during each interruption. As can be seen from the

data, the time required to shut the electronic componentof the cell off was on the

order of 50 microseconds. Such a rapid shut-down rate should be construed to be an

advantage from an operating (safety) point of view.

In addition to proceeding from a power producing modeto open circuit, the response

times were also monitored for whenthe system is at open circuit and then dropped

across a resistance to generate the 100 amps/ft (. The results of these tests are

essentially identical to those whenthe current path was opened, both in behavior as

well as response time.

For a power device to be practical the cell response time must be fast and efficient

and without the generation of transients. In both of the above response time tests,

the time required to reach various percentages of full power are presented in Table

3. In all cases, it is seen that the time to reach full "on" power was short, but in

all cases 90%of full power was achieved within 80 milliseconds.

Finally, since the fuel cell was turned off and on several times throughout the test,

it is important to demonstrate the reliability and resiliency of the PEM system by

presenting polarization data after such system upsets. The series of polarization

curves generated throughout the week-long run was carried out in order to determine

if changes were occurring to the membrane/electrode assembly. Presented in Figure 10

is a series of current-voltage sweeps(IR included) taken at the times labeled on

each curve. After the numerouson-off cycles, as determined from the polarization

behavior, no apparent degradation was taking place.

117
CONCLUSIONS

The results of tests utilizing a new experimental membrane useful in proton-exchange

membrane fuel cells were presented. The high voltage at low current densities can

lead to higher systems efficiencies while, at the same time, not sacrificing other

Critical properties pertinent to membrane fuel cell operation. A sbrie_ of tests to

determine response times indicated that "on-off" cycles are on the order _f 80 milliseconds to reach 90% of full power. The IR free voltage at 100 amps/ft = was determined _nd the results _ndicating a membrane/electrode package resistance to be .15

ohm-cm _ at 100 amps/ft =.

REFERENCES

I. McElroy, J., General Electric Co., Solid Polymer Electrolyte Fuel Cell Technology

Program, Final Report. NASA Contract # Nas 9-15286, May, ]g80, p 5.

2. Ezzell, B. R.; Carl, B.; Mod, W. A.; "Ion Exchange Membranes for the Chlor-Alkali

Industry," AIChE Symposium Series, Houston, Texas, March, 1985, pp 49-51.

, Eisman, G. A., "The Physical and Mechanical Properties _f a New Perfluorosuifonic Acid Ionomer for use as a Separator/Membrane _n Proton Exchange

Processes," Proceedings of the 136Th Electrochemical Society Meeting, Boston,

May, 1986.

4. Eisman, G. A., " The Application of a New Perfluorosulfonic Acid Ionomer in Proton-Exchange Membrane Fuel Cells: New Uitra-HighCurrent Density Capabilities,"

Proceedings of the International Seminar on Fuel Cell Technology and Applications, The Netherlands, 26 October, 1987, pp 287-294.

5. Watkins, D., canadian Fuel Cell:Program, International Seminar on Fuel Cell Technology and Applications, The Netherland s , 26 October, 1987.

6. Vanderborgh, N. E.; Spirio, C. A.; Huff, J. R.; "Advanced Fuel Cell Systems for

Transportation Applications," Proceedings of the international Seminar on Fuel

Cell Technology and Applications, The Netherlands, 26 October, 1987, pp 253-262.

= :

7. Hoadley, J., Lawrence, R., "Scale-up of a Proton-Exchange Membrane Fuel Cell,"

Proceedings of the Fuel Cell Seminar, Long Beach, CA, October, 1988, pp 263-266.

|

=

_e_

z

=

B

|

R

i

118

i

r
FIGURg !

¥OLTAG]I-CURRgNT POIARIYATION DATA

!

0.9

0.8

0.7

0.6

0.5

0.4

0.3

0.2

0 500 t000 t500 2000 2500 3000 3500

CURRENT DENSITY (ASF)

HP./02 40/8O PSZGBOC

2 MG/Q42/SIDE Pt

DO#

NAFION-it7

BO PSI-02/40 PSI"H2

..,...,........

FIGURE 2

CHEMISTRY OF THE DOW CHEMICAL COMPANY

A

0 0 0

FSO=CF,CF + CICF:CFCF= .... > FSO=CF=CF=OCF_F

(_F=CI

A + Na2CO_ ..... >

B

FSO,CF,CF2OCF=CF,

B + CF_=CF_ .... > POLYMER

119
FIGURE 3

POLYMER STRUCTURES

±

D0W

-CF2CF'2cF1CF2CF2-

1

O

I

OF2

t

OF'2

i

O--S'--O

i

OH+

DUPONT NAFION

-CF2CF2CF2CF1 CF2CF2-

6

I

CF2

I

CF2

I

0

I

CF2

I

CF2

I

O---S--O

I

OH-IFIGURE4

CELLHARDWARE

®

IL, (_BUSS

!

FLUID

ENDPLATE

120

14&E

LEAVES m---l_

PNEUMATVC

ENOPLATE
FIGURB 5

VOLTAGg - CURRBNT POLARIZATION DATA

HYDROGXN OXYGEN 80 C

ORIGINAL PAGE IS

OF POOR QUALITY

1t1¢I',MII

0.9

o.7°'e_ :k _ :k _

0.6 i

i

0.5

0.41

0.3

0.2

0.i

0

0

4O16Oi_I!1 141/02

i. t 91"0101H_ 2 Sl"OICH02

_0 Pt/CH'2/S,T_

_o 5oo 715o ,ooo _o _Boo _o 2ooo _5o

0.9

0.8

0.7

0.6

0.5

0.4

0.3

0

FIGURB 6

VOLTAGB-CURRgNT POLARIZATION DATA

0 --> 500 A_

:k :k

lO0 200 300 400 500

commm' :mm'rr Ump)

4O/6O FSIS H2/02 eo c

l.t 5TOZCH1t2:2 STOICH02

eeo Pt/CH'e/S_E

121
FIGURE 7

YOLTAGE VS. TIME (HRS)

200 ASF

ORIGINAL P_GV. _S

OF POOR QUALITY

TOLTAGll

0.80"gf_ _k >_ * _ _

°70.6 f

0.5_

0.4!

[

0.3 Lo 10 20 30 40 50 60 70 80 90 t00

40/60 PSIG H2/02

t.t 5TOiCH H_ 2 9TOICH 02

2mg Pt/CH'2/SiDE

80 C

122
IqGUlW 9

INT]D_UPTION

100 _ --> OCV

ORIGIP,]AE PACE |3

OF POOR QU:&iTY

0.g4

0.93

0.92

0 50 i00 i50 200 250 300 350 400

4O/6O !=916 H2/02 8O C

i.t 91"0ICHH_ 2 91"0ICH02

_iCN'Z/SIOE

.,L.._-= _ -,.1 _-a.._ --.1 -- -a----.I

45O 500 55O 6OO

200

IelGURI l0

VOLTAGI VS. C_ DXNSITY

POLk2IZ&TION DATA VS. BUN

|

3OO 400 500 500 700 BOO 900

123

TIME - INITIAL

TIME - 5 HRS.

0

TIE - iO t4:lS.

TIME - 25 H_.

TIME - 50 HRS.

--$--

J

iO00
TABLE 1

MEMBRANE PROPERTIES

tONiC CONDUCTIVITY .1 - .2 (0wU-CM),,.-1

TBNSILB STRI_GTH 2500 - 4000 Pm

WJLT3_ UPTAJ[B 58 - 60_

vmm Pmmsarr¥ +.5 • s-_ cc/(ci_*ssc)

Gag PRRM_TION

m_soGz_ 5 x z-9 CC-CM/(CM--_*CM_z*ssc)

OXYGI_q 2.5 X B-9

TABLE 2

IR FREE VOLTAGE VS. CURRENT DENSITY

CUaaeNTDs_rsr_ £J-_l 9z____-_cx__ __vo. LT,¢S(MVJ

z

=

25

5o

75

100

3

6

11

15

124
TABLE 3

RESPONSE TIME VS. PERCENT FULL POWER

OPEN CIRCUIT --> 100 ASF

FULL POWER TIME (MS)

10 2

2O 5

30 8

40 12

50 16

60 21

70 28

80 41

90 51

97 81

OPEN CIRCUIT VOLTAGE 1 .O23 V

VOLTAGE AT 100 ASF" .915 V

"125

N90-20467

HYDROGEN-OXYGEN PROTON-EXCHANGE MEMBRANE FUEL CELLS

AND ELECTROLYZERS

R. Baldwin

NASA Lewis Research Center

Cleveland, Ohio 44135

M. Pham

NASA Lyndon B. Johnson Space Center

Houston, Texas 77058

A. Leonida, J. McEIroy and T. Nalette

Hamilton Standard

Windsor Locks, Connecticut 06096

INTRODUCTION

Hydrogen-oxygen SPE fuel cells and SPE electrolyzers (products of Hamilton

Standard) both use a Proton-Exchange Membrane (PEM) as the sole electrolyte.

These solid electrolyte devices have been under continuous development for

over 30 years. This experience has resulted in a demonstrated ten-year SPE

cell life capability under load conditions.

Ultimate life of PEM fuel cells and electrolyzers is primarily related to the

chemical stability of the membrane. For perfluorocarbon proton exchange

membranes an accurate measure of the membrane stability is the fluoride loss

rate. Millions of cell hours have contributed to establishing a relationship

between fluoride loss rates and average expected ultimate cell life. Figure 1

shows this relationship. Several features have been introduced into SPE fuel

cells and SPE electrolyzers such that applications requiring Z 100,000 hours

of life can be considered.

Equally important as the ultimate life is the voltage stability of hydrogenoxygen fuel cells and electrolyzers. Here again the features of SPE fuel

cells and SPE electrolyzers have shown a cell voltage stability in the order

of 1 microvolt per hour. That level of stability has been demonstrated for

tens of thousands of hours in SPE fuel cells at up to 500 amps per square foot

(ASF) current density. SPE electrolyzers have demonstrated the same stability

at 1000 ASF.

Many future extraterrestrial applications for fuel cells require that they be

self recharged (i.e., regenerative fuel cells). This requirement means that a

dedicated fuel cell and a dedicated electrolyzer work in tandem as an

electrical energy storage system. Some applications may find advantages with

a unitized regenerative fuel cell (i.e., one cell that operates alternately as

a fuel cell and as an electrolyzer). Electrical energy storage for earth

orbits via hydrogen-oxygen regenerative fuel cell systems can have specific

energies in excess of 50 watt-hours/kg (Ref. I). For extraterrestrial surface

electrical energy storage the hydrogen-oxygen regenerative fuel cell can have

significantly increased specific energies:

Mars Base - 500 Watt-hours/kg

Lunar Base - 1000 Watt-hours/kg (Ref. 2)

127

PRECEDING PAGE BLANK NOT FILMED
To translate the proven SPE cell life and stability into a highly reliable

extraterrestrial electrical energy storage system, a simplification of

supporting equipment is required. Static phase separation, static fluid

transport and static thermal control will be most useful in producing required

system reliability. Although some 200,000 SPE fuel ceil hours have been

recorded in earth orbit with static fluid phase separation, no SPE

electrolyzer has, as yet, operate d in space.

Under NASA sponsorship a flight experiment of a unitized regenerative fuel

ceil is being studied. If selected for actual flight under the NASA 0AST

Outreach Project, several advanced features will be tested in space.

The objective of the flight experiment is to test the space viability of the

incorporated features, and not to imply that the specific configuration of the

flight experiment is optimum for any given extraterrestrial application. With

a successful flight experiment, supported by terrestrial experiments, the

system designer can select the proven advanced system features that are

appropriate for any particular extraterrestrial application.

ULTIMATE CELL LIFE

The ultimate life of PEM fuel cells and eIectrolyzers, based on fluoride loss

rate, assumes that the cell, stack and system designs are configured to

prevent premature mechanical failures. With the elimination of mechanically

induced failure, the fluoride loss rate has been found to be an excellent

measure of the health and life expectancy of PEM cells.

The fluoride lost from the PEM cell is actually a degradation product of the

perfluorocarbon membrane. It should, therefore, not be surprising to find the

fluoride loss/ultimate life relationship. The Figure I relationship was

primarily based on PEM cells using perfluorocarbon membranes in the 8 to 12

mils thickness range. The Figure suggests that perfluorocarbon membranes of

less than 8 mils will have a similar ultimate life with a proportional loss of

fluoride. This, of course, will only be verified when a large number of cells

containing advanced thin perfluorocarbon membranes have accumulated

substantial lifetimes.

The fluoride loss rate determination for non-evaporatively cooled PEM fuel

cells can be directly measured in the liquid product water. Fluoride

detection at the low ppb range can be made directly, while detection at the

ppt range is accomplished by concentration techniques. To put the Figure i

curve into everyday calibration, the following PEM fuel cell example is

offered:

Cell CharacteristiG_

Membrane Thickness - 10 mils

Current Density 500 Amps per Square Foot (ASF)

128
Life Prediction Vs. Fluoride in Product H2_

1,000 hours -

10,000 hours -

100,000 hours -

2 ppm

200 ppb

20 ppb

The fluoride loss rate determination for liquid anode feed PEM electrolyzers

can be directly measured in the open loop "proton" water. The following PEM

electrolyzer example is offered:

Cell Characteristics

Membrane Thickness - 10 mils

Current Density - 1000 ASF

Life Prediction Vs. Fluoride in Proton H2_

1,000 hours - 140 ppb

10,000 hours - 14 ppb

100,000 hours - 1.4 ppb

Since long life of PEM cells has been inversely equated with fluoride loss

rate, the designer can quickly determine the suitability of a given

configuration for a particular application. Each manufacturer of PEM cells

undoubtedly has his own proprietary techniques for the control of fluoride

release rate.

SPE fuel cell and electrolyzer designers incorporate proprietary features into

system configurations for extended life. One such demonstration SPE cell is

within seven months of reaching 100,000 hours of high current density

operation. Figure 2 displays this cell as it passed the ten year mark in May

1988. Two additional SPE test cells have accumulated over ten years of

"uninterrupted" operation since the May 1988 milestone.

CELL VOLTAGE STABILITY

The voltage stability of PEM fuel cells and electrolyzers is a second very

important feature for many applications. Experience has shown that features

that provide a long cell life, as indicated by low fluoride release rates, do

not necessarily equate long cell life with stable voltage. This is indeed

unfortunate, as it makes it next to impossible to predict the long term

voltage stability of a cell design from short-term tests.

In earlier SPE cell designs a reasonable voltage-change rate would

unexpectedly turn to a high voltage-change rate. When unpredicted mechanisms

were identified and corrective measures taken, the level and longevity of

voltage stability improved. As one might imagine, the test, analysis,

corrective action and retest cycles have taken many years to obtain the

current level of voltage stability.

Figure 3 shows the voltage stability, as measured by voltage-change with time,

of six SPE test units. The baseline configuration for SPE electrolyzers and

SPE fuel cells is represented by SPE electrolyzer C and SPE fuel cell E. Both

of these units have a voltage-change rate of 1 microvolt/cell hour at 1000 ASF

and 500 ASF, respectively.

129
SPEelectrolyzers A and B purposely deviate from the baseline configuration

for application specific reasons. Both of these units represent a voltagechange rate of 7 microvolts/cell hour. This voltage-change rate is perfectly

acceptable for the intended applications (i.e., electric utility hydrogen

generation and nuclear submarine oxygen generation).

Like the SPE fuel cell baseline E, SPE fuel cell D displayed a voltage'change

rate of 1 microvolt/cell hour. The unit D configuration, formally considered

the SPE fuel cell baseline, is limited in current density. SPE fuel cell F

had a voltage-change rate of 6 microvolts/cell hour and represents an earlier

configuration which was part of the development data base leading to the

present SPE fuel cell baseline.

1

The configurations of the baseline SPE electrolyzer and SPE fuel cell consist

of a combination of patented and proprietary features. Each PEM electrolyzer

and PEM fuel cell manufacturer will have their own features which, like the

SPE devices, will have to be verified by the life test process.

EXTRATERRESTRIAL APPLICATIONS

Many of the future extraterrestrial applications for hydrogen-oxygen fuel

cells and electrolyzers involve combining the two devices to produce the

equivalent of a battery for electrical energy storage (see Figure 4). The so

called regenerative fuel cell has the distinct advantage over traditfonal

batteries in that power and energy are separated. One needs only to increase

reactant storage, without increasing the reactor stack(s), to increase the

stored energy. The results of this distinction is that the mass advantage of

typical advanced batteries lose out to regenerative fuel cells when the

discharge time is increased beyond a few tens-of-minutes.

Energy storage system sizing has been examined for a variety of extraterrestrial applications. The most favorable application _ of the

hydrogeh,oxygen regenerative fuel cell is the lunar base application because

of the long charge/discharge time span. The studies have indicated the next

best available energy storage technology (NaS battery) is more than ten times

as massive as the hydrogen-oxygen regenerative fuel cell system.

The state-of-the-art of the hydrogen-oxygen regenerative fuel cell is not

without disadvantages. Even considering the demonstrated lifetime and voltage

stability of the SPE devices, the questions of efficiency and complexity must

be considered.

The energy-out-t0-energy-in ratio range generally accepted for state-of-theart hydrogen,oxygen regenerative fuel cells is 55% to 65%. Advanced batteries

operating at 75% energy storage efficiency or greater have the clear advantage

in this characteristic. The primary reason for the lower hydrogen-oxygen

regenerative fuel cell energy storage efficiency is related tO electrode

performance. Years of research have been invested to improve low-temperature

electrode performance with only minor increases observed. Without a

breakthrough in low-temperature electrode performance,the lower efficiency of

the state-of-the-a_t low-temperature hydrogen-oxygen regenerative fuel cell is

fundamental.

130
The efficiency of the hydrogen-oxygen regenerative fuel cell results in a

greater energy input requirement for a fixed output when compared to advanced

batteries. In the case of a photovoltaic power source, this means an increase

in the size of the solar array. However, as charge times increase this

differential mass becomes insignificant. An order-of-magnitude lower mass for

the overall lunar base power system based on the hydrogen-oxygen regenerative

fuel cell, as compared to advanced batteries, is still projected even when

including the state-of-the-art photovoltaics.

Although the efficiency disadvantage of the hydrogen-oxygen regenerative fuel

cell turns out to be insignificant in most extraterrestrial applications, the

complexity of state-of-the-art regenerative fuel cells is much greater than

most advanced batteries. The very feature that gives the hydrogen-oxygen

regenerative fuel cell its significant mass advantage is responsible for its

complexity. The complexity arises primarily due to the storage of reactants

and products external to the reactors. These state-of-the-art complexities in

hydrogen-oxygen regenerative fuel cells include:

Fuel cell product water removal from the reactant chamber by dynamic

reactant recirculation.

Dynamic separation of the fuel cell product water from the reactant.

Pumping electrolysis process water up to gas generation pressure.

Temperature-controlled fluid recirculation in both fuel cells and

electrolyzers.

Separate dedicated fuel cells and electrolyzers.

In total, the state-of-the-art hydrogen-oxygen regenerative fuel cell system

would use in the range of 4 to 6 rotating devices, not including redundancy,

to provide fluid transport and/or phase separation. Additionally, on the

order of twice as many individual cells would be required as would be needed

for advanced batteries, simply because the batteries operate both in charge

and discharge modes. The complexity, in terms of dynamic components and the

large piece-part count, with double the number of cells, relegates the

state-of-the-art hydrogen-oxygen regenerative fuel cell to a lower reliability

than the more massive batteries.

The reliability of the hydrogen-oxygen regenerative fuel cell can be increased

by component redundancies. This approach diminishes the mass advantage and

adds to the system complexity. Our approach, of a unitized (reversible)

hydrogen-oxygen regenerative fuel cell with static fluid and thermal control

features, can improve reliability and reduce complexity, which further

improves the mass advantage over advanced batteries.

SYSTEM SIMPLIFICATIONS

With complexity and lower reliability identified as the major concerns in

hydrogen-oxygen regenerative fuel cells, manufacturers have begun to respond

with a variety of approaches. In particular, SPE fuel cell and SPE

electrolyzer designs are being directed toward reduced complexity. Specific

developments include:

w

i

m

m

Reversible SPE cell

Passive fuel cell product water separation and removal

Passive electrolyzer process water introduction

Passive electrochemical gas pumping

Passive heat management

131
Separate dedicated SPE fuel cells and SPE electrolyzers came into favor

several years ago due to higher overall energy storage efficiency. It is true

that in reversible SPEcells, the preferred SPEfuel cell catalyst results in

poor electrolyzer performance and the preferred SPE electrolyzer catalyst

results in poor SPE fuel cell performance. Compromise catalyst systems have

been developed which bring the overall energy storage efficiency within 5% of

the dedicated configuration. Figure 5 shows performance curves for these

various configurations.

A specific application may find the 5% efficiency reduction as an acceptable

trade for simplicity and mass reduction. Alternately, if the application

dictates, by trade study, that discharge efficiency is of prime importance, a

10% energy storage efficiency reduction may be preferred while using optimized

fuel cell catalyst.

Passive fuel cell product water separation and removal from the reactor stack

in microgravity is certainly not a new idea. Both the Gemini and Biosatellite

SPE fuel cells operated with passive wicking product water systems. These

systems, however, operated only at low current densities with higher mass and

volume structures.

Recent developments in passive water separation and removal from SPE fuel

cells have involved separation of the liquid product water from the gaseous

oxygen within the cell structure (Ref. 3). This approach, shown in Figure 6,

eliminates the need for wicking and uses a hydrophilic separator plate in the

oxygen cavity. Testing of this configuration has been very successful to date

with operation demonstrated against gravity (i.e.,- product water separated

upward from the cathode electrode). Figure 7 displays the performance of an

SPE fuel cell with the passive product water removal feature operating against

gravity.

The passive water vapor feed SPE electrolyzer developed during the seventies

used a hydrated ion-exchange membrane to feed water vapor to the operating

cell. This cell structure (shown in Figure 8) eliminated the need for

separating liquid water from the products. However, close control of the

water pressure was required. If the water pressure was significantly higher

than hydrogen pressure, water hydraulically passed through the water feed

barrier and exceeded the electrolysis rate. Conversly, if the water pressure

was below the hydrogen pressure, an accumulation of hydrogen in the water feed

chamber by normal diffusion resulted in water starvation.

Recently, an inovative feature has been incorporated into the passive water

vapor feed SPE electrolyzer. This feature allows the water pressure to be

well below hydrogen pressure without the concern of hydrogen buildup in the

water feed chamber. The feature consists of electrifying the water feed

membrane in order for it to function as an electrochemical hydrogen pump (Ref.

4). Figure 9 shows the electrification of the water feed membrane. The

fraction of a watt required for each cell is considered a small penalty to pay

for the elimination of the requirements for pumping water up to pressure and

maintaining close pressure control.

The electrochemical hydrogen pump allows the water to be fed to the cell at

any source pressure down to the cell water vapor pressure without any

mechanical pumping. Figure 10 shows the water vapor feed SPE electrolyzer

performance with water pressure 160 psi below hydrogen generation pressure.

132

Z
Several thousand cell hours have now been accumulated on electrolyzer cells

supported by the use of electrochemical hydrogen pumps.

The simplifying features described herein have all demonstrated successful

operation in ground testing. Although passive fluid phase separation was

successfully used in the Gemini and Biosatellite SPE fuel cell systems, none

of the described simplifying features have actually flown in space.

THE FLIGHT EXPERIMENT

The planned unitized regenerative SPE fuel cell experiment makes use of the

proven life and voltage stability features as well as the simplifying aspects.

A basic fluid schematic of the flight experiment is shown in Figure 11. This

figure shows one of several experiment options studied. If selected, the

experiment will fly on a future Shuttle mission in a simulated low earth

orbit energy storage configuration. The potential one week experiment should

be of ample length to assess the performance of the various system

simplification aspects.

During the experiment discharge cycle, hydrogen gas and oxygen gas are

delivered to the reversible cell on a purely demand basis. The reversible

cell converts these reactants into dc power with a current output of

approximately i00 ASF. The product water, formed within the oxygen chamber as

a liquid, contacts the porous plate. The spring bellows in the oxygen/water

storage tank creates a pressure differential across the porous plate (0_ >

HpO) of about i psi. The i psi is more than sufficient to force the product

w_ter through the hydrophilic porous plate at the generation rate, but, not

nearly high enough to allow passage of oxygen gas.

The product water will be saturated with oxygen, and some small bubbles of

oxygen gas will appear as the water pressure decreases by 1 psi as it passes

through the porous plate; some oxygen will come out of solution. As this

water passes through the water chamber of the water vapor feed barrier

membrane, the oxygen gas and dissolved oxygen are converted to water at the

positive electrode of the hydrogen pump. This oxygen is chemically consumed

by combining with the diffusing hydrogen that otherwise is electrochemically

pumped back to the hydrogen chamber. The water eventually delivered to

storage is essentially free of dissolved gases. This is an important feature

as the system pressures alternate up and down through the charge/discharge

cycles.

At the end of the discharge period the pressures in the reactant storage tanks

are at the reduced level and a load contactor is opened. This is followed by

closing a power contactor to supply charge power to the reversible cell.

In the electrolysis charge mode, water vapor is fed to the operating

electrolysis cell across the hydrogen compartment gap. Hydrogen and oxygen

gases are delivered to their respective tankage and the system pressures rise.

An accumulation of water in the reactant storage tankage is prevented by

thermally matching the tankage with the operating cell and the fact that the

produced gases are in fact subsaturated.

133
The electrochemical hydrogen pumpmilliwatt power draw continues through both

the charge and discharge modes. This assures that hydrogen gas does not block

the flow of water, and it also provides an actual measure of the amount of

oxygen consumedduring the discharge water desaturation step.

The entire experiment will be cold-plate mounted. Heat generated within the

cell hardware is conducted to the base plate. Figure 12 is a photograph of

the experimental hardware mock-up.

_UMMARY

In summary, a flight experiment is planned for the validation in a

microgravity environment of several ground proven simplification features

relating to SPE fuel cells and SPE electrolyzers. With a successful

experiment these features can be incorporated into the designs of equipment

for specific extraterrestrial energy storage applications.

I °

.

3.

.

REFERENCES

Appleby, J.: Regenerative Fuel Cells for Space Applications.

Journal of Power Sources, 22 (1988) p.p. 337-385.

NASA Conference Publication 3016: September 12-13, 1988 p.p. 206

Leonida, A: Hydrogen/Oxygen Fuel Cells with In-Situ Product Water

Removal. 33rd International Power Sources Symposium, June 13-16, 1988

Sedlak, J; Austin, F.; and LaConti, A: Hydrogen Recovery and Purification Using the Solid Polymer Electrolyte Electrolysis Cell,

International Journal of Hydrogen Energy, (1981) Volume 6, p.p. 45-51

134
ORIGINAL PAGE IS

OF POOR QUALITY

,GO.dOG

lO.(JO0

1000

\

100

10"? 1043

FLUORIDE LOSS RATE --

10"5 10 .4

GRAMS

FTI.MIL.HR

tO'3

I? ELEC,R,C SUE T I

12 i

r ' !

' _ t r-"+

! ST_*GE MOOULE,'S,

I "_--" l _ PLU,OL,NE

SOLAF t l --- _ ELECTR)CAL UNE

ARRA_ I 1

I RADIATOR

FIGURE 1. EXPECTED LIFETIMES OF SPE CELLS WITH PERFLUOROCARBON

PROTON EXCHANGE MEMBRANES

_21:I::=_:_Z_:_I?I_  ̧:::::I:i:_i_i::

: __-:_:_::---_GtRO Ly S]_3 :(--7:2::;

FIGURE 4. HYDROGEN-OXYGEN REGENERATIVE FUEL CELL ENERGY STORAGE SYSTEM

• STANbARI_ IrC pT]_T -

E_ ORU _CUAVE_

OD.g__ STA_D COMPROMISE E[ECT_LyS_S CATALYST E mT'E §_E 5_0 I

I

/

CUR_ O_NSlT_ <ASF_

FIGURE 2. "UNINTERRUPTED" SPE ELECTROLYZER LIFE TEST

FIGURE 5. COMPARISON OF PERFORMANCE DURING FUEL CELL AND ELECTROLYSIS MODES

2,

20

1,9

> f8

I

g 7

• o_

> oe_

F

_ 05

_ O4

03

01

C

SPE ELECTROLYZER BASELINE

D

E

ERE FUEL CELL BASEL!NE

SPE ELECTROLYZERS

A -- 10 CELL STACK (2000 ASF)

B -- SINGLE CELL (1000 ASE)

C -- SINGLE CELL (1003 ASF)

10,000 20,OOO

HOUR5 AT STATED CURRENT DENSITY

i

O -- 4 CELL STACK {120 ASFI

E -- SINGLE CELL (500 ASF)

F - SINGLE CELL (500 ASE)

1

30 ,(1_ 10.1100

POROUS PLATE

SEPARATOR

CATHODE JlIi o ----=.- I

FRAME

I H20 __ [ '---I_m_'J _ GASKET

_ _ _1-_--_ CATHODE FRAME

-..--- .... . _ _ MEMBRANE/ELECTRODE

:_ _---'-- - -'_ _ __--_]'_'_ ASSEMBLY

H2 r_._ _l ._1_ ANODE FRAME

,, ,

GASKET

I

ANODE -- Q _ I

I

FIGURE 3. SPE CELL VOLTAGE STABILI13" FIGURE 6, PASSIVE STATIC WATER REMOVAL SPE FUEL CELL -- CROSS SECTION

:. ::-.;.._

]35
t0

0,9

<J 0,8

07

.d

8

0._

0.0

CELL AREA: 0_23 FT _

MEMBRANE: NAFION 120

TEMPERATURE: 150"F

! l ,,,L, I I I I

50 100 150 A= F _ 200 250 300 350

I I I T I I I I

10 20 30 40 50 60 70 80

CURRENT OUTPUT_ A

FIGURE 7. PASSIVE WATER REMOVAL 5PE FUEL CELL

PERFORMANCE AT I_"F

I °-- i

I _ GASKET

H20 __-..=e_- FEED WATER FRAME

1 __ -- .... ''-- -- _ WATER FEED BARRIER

__ H 2 _'_ -,,,.=l-.---,-. CATHODE FRAME

[ r -_-_:_ -_ __.___-- - ___= ._ .... --.1 MEMSRANEIELECTRODE 1_e_'- ASSEMBLY

"1- 02 __,_1.,,._'_ ANODE FRAME

I I---_ASKET l o-I

FIGURE 8. WATER VAPOR FEED SPE ELECTROLYZER

I o____J

[ _ GASKET

&%%%o_ ._o _ FEE° WATER FRAME

f.---- _ -_ _ ._ . __ - _ WATER FEED BARRIER

_ ___'c--" .... -- _ _ _ "Jl'-'_ WITH H 2 PUMP

-- H 2 _ _",_'--_ CATHODE FRAME

[.._::______--__.--_---__,gg_#;_/,ELECTRODE

Jr 02 __]_ ANODE FRAME

_ GASKET I o-J

FIGURE g. WATER VAPOR FEED sPE ELECTROLYZER WITH ELECTROCHEMICAL

HYDROGEN PUMP

2.0

L)

a

> 1.9--

I

g_

> 17

d

0.O

ORIGINAL PAGE IS

OF POOR QUALITY

FEEDWATERJ PRESS i-- i35 PSIA I

r'] TEMP 1BO°F

0 10 20 30 40 50 60 70

I CELL CURRENT (AMPS) I I 1 I I I

O 50 1 O0 150 200 250 30'0

CURRENT DENSITY (ASF)

I

350

FIGURE 10. WATER VAPOR FEED SPE ELECTROL_ER PERFORMANCE

oz + H_

t" .... RS ........ I I'

LECTROCHEMICA L PUM

H20

---'p=" O

Vl

H20

FIGURE |lr UNITIZED REGENERATIVE SPE FUEL CELl.

z

FIGURE t2. UNmZED REGENERATIVE SPE FUEL CELL MOCK.UP

i

=

i

=i L

z

r

z

m

136

ORIGINAL- PAGE'

BLACK AND WHITE PHOTOGRAPH
N90-20468

ELECTROCATALYSIS FOR OXYGEN ELECTRODES IN FUEL CELLS AND

WATER ELECTROLYZERS FOR SPACE APPLICATIONS

Jai Prakash, Donald Tryk and Ernest Yeager

Case Center for Electrochemical Sciences

and the Chemistry Department

Case Western Reserve University

Cleveland, Ohio 44106

The lead ruthenate pyrochlore Pb2Ru206.5 in both high and low area forms has

been characterized using thermogravimetric analysis, X-ray photoelectron spectroscopy, X-ray diffraction, cyclic voltammetry and 02 reduction and generation

kinetic-mechanistic studies. Mechanisms are proposed. Compounds in which part of

the Ru is substituted with Ir have also been prepared. They exhibit somewhat

better performance for 02 reduction in porous gas-fed electrodes than the unsubstituted compound. The anodic corrosion resistance of pyrochlore-based porous

electrodes was improved by using two different anionically conducting polymer

overlayers, which slow down the diffusion of ruthenate and plumbate out of the

electrode. The 02 generation performance was improved with both types of electrodes. With a hydrogel overlayer, the 02 reduction performance was also improved.

INTRODUCTION

In most instances separate electrocatalysts are needed to promote the reduction

of 02 in the fuel cell mode and to generate 02 in the energy storage- water

electrolysis mode in aqueous electrochemical systems operating at low and moderate

temperatures ( T < 200 °C). This situation arises because, even with relatively

high performance catalysts, the 02 reduction and generation reactions are still

quite irreversible with much overpotential. The potentials of the 02 electrode in

the cathodic and anodic modes are separated by typically 0.6 V and the states of the

catalyst surface are very different. Interesting exceptions are the lead and bismuth ruthenate pyrochlores in alkaline electrolytes. These catalysts have high

catalytic activity for both 02 reduction and generation (1-6). Furthermore,

rotating ring-disk electrode measurements provide evidence that the 02 reduction

proceeds by a parallel four-electron pathway (6). The ruthenates can also be used

as self-supported catalysts to avoid the problems associated with carbon oxidation,

bu.t the electrode performance so far achieved in our laboratory is considerably

less.

At the potential involved in the anodic mode the ruthenate pyrochlores have

substantial equilibrium solubility in concentrated alkaline electrolyte. This

results in the loss of catalyst into the bulk solution and a decline in catalytic

*This work was supported by the NASA-Lewis Research Center under Contract No.

NAG3-964.

137
activity. Furthermore, the hydrogen generation counter electrode may become

contaminated with reduction products from the pyrochlores (lead, ruthenium).

A possible approach to this problem is to immobilize the pyrochlore catalyst

within an ionically conductive solid polymer, which would replace the fluid electrolyte within the porous gas diffusion 02 electrode. For bulk alkaline electrolytes,

an anion exchange polymer is neede@with a transference number close to unity for

the OH- ion. Such a membrane may not block completely the transport of the lead and

ruthenium, which are expected to be in complex anionic forms. Preliminary short

term measurements with lead ruthenates using either i) a_partially fluorinated anion

exchange membrane as an overlayer on the porous gas-fed electrode or 2) a hydrogel

coating on the porous electrode indicate lower anodic polarization and similar

cathodic polarization. With the hydrogel coating, there is also some enhancement of

the performance at higher current densities.

EXPERIMENTAL

The pyrochlores were synthesized at CWRU using the alkaline solution technique

of Horowitz et al (3). Samples of Pb2Ru206. 5 and PD2[RUl.67Pb0133]06. 5 were also

provided by Exxon. This method involves the reaction of the appropriate metal

cations by precipitation and subsequent crystallization of the precipitate in a

liquid alkaline medium (4 M KOH) in the presence of 02 at ~75-90°C. The salts used

were Ru(NO)(N03)3, in aqueous solution, 1.5 _ (w/v) from Strem, IrCI3.3H20

(Aldrich, Gold label) and Pb(CH3COO)2-3H20 (MCB, reagent grade). The salts were

added in the appropriate amounts to achieve the desired stoichiometries. The

reactions were carried out for 24-72 h, until the supernatant solution was free of

detectable amounts of the reactant metal ions. The suspension was filtered and

washed with water and glacial acetic acid. The X-ray diffraction (XRD) was measured

after a heat treatment (HT) in air at 300-350°C for 12 h. In some cases XRD was

also run prior to the HT.

Thermogravimetric analysis was performed on a Perkin-Elmer 7 Series Thermal

Analysis System With O2 and N 2 purges at a temperature ramp rate of 15°C min -I Xray photoelectron spectroscopy was obtained using a Varian lEE- 15 instrument.

A modification of the ring-disk electrode (RRDE) technique was used in which

a thin PTFE-bonded porous coating of the pyrochlore was applied to a basal plane

pyrolytic graphite disk, which was slightly recessed_ PTFE suspensi0n (Tefion T30B,

Du Pont) was diluted with water and ultrasonically agitated with the pyrochlore and

filtered with a I #m pore size polycarbonate filter membrane (Nucleopore). The

resulting mixture was 5 wt _ in PTFE solids. The semi-moist _aste was kneaded with a

spatula and then applied to the disk (either 0.196 or 0.45 cm L) and spread evenly.

Excess material was carefully removed from the edge of the disk. The layer was not

allowed to dry in order to avoid cracking of the layer and was placed in the

electrolyte solution. The amount of pyrochlore in the layer was determined after

the experiment by fir§t washing it in water_ removing it, drying and Weighing. The

typical loading was 4-6 mg cm 2. The Au ring had a collection efficiency N of 0.177.

The Teflon electrochemical cell had separate compartments for the working RDE or

RRDE electrode, counter electrode (Au) and reference electrode (Hg/HgO, i M OH-).

The KOH solution was prepared from the solid (Fluka, puriss., p.a.) by dilution

with triply distilled reverse osmosis water.

138
Porous gas-fed electrodes were fabricated as follows. Teflon T30B suspension

was diluted approximately to 2 mg cm "3 in water and slowly added to an aqueous

suspension of 38 mg of pyrochlore and 35 mg of air-oxidized (-I h at 600°C)

Shawinigan black (Chevron Chemical Co., Olefins and Derivatives Div., Houston,

Texas). The suspension was filtered using a l#m pore size filter membrane. The

resulting paste was kneaded with a spatula until slightly rubbery, shaped into a

1.75 cm diameter disk in a stainless steel die using hand pressure and then pressed

at a pressure of -i00 kg cm "2. The disk was applied to a 0.5 mm thick disk of

Teflon-carbon black hydrophobic backing material containing silver-plated nickel

mesh (Electromedia Corp., Englewood, NJ). A final pressing was done at -300 kg

cm -2 at room temperature and then the disk was heat-treated at -280°C for 2 h in

flowing helium.

In some cases a 1.75-cm diameter disk of a partially fluorinated anion exchange

membrane material (Type 4035, RAI, Haup_age, NY) was pressed onto the electrolyte

side of the electrode at 200 kg cm -L. In other experiments a mixture of

poly(dimethyldiallylammonium)chloride (DMDAAC, 15 wt _ in aqueous solution, Polysciences) and Nafion 117 (5 wt_ in alcohol solution, Aldrich) was diluted by a

factor of 6 by volume and painted onto the electrolyte side of the porous electrode

and allowed to air-dry. The resulting hydrogel film was -1.5 mg cm -2

The porous gas-fed electrode was placed in a special screw-cap-type

KeI-F holder which holds the electrode vertically between a Pt foil current

collector on the gas-side and a polyethylene-polypropylene rubber gasket on the

electrolyte side. A stream of either air or 02 was directed at the back of the

electrode. The area exposed to the electrolyte was 0.97 cm 2. The 02 reduction

measurements were performed galvanostatically in 5.5 M KOH at 25°C with a fastaction potentiostat (BC-1200, Stonehart Associates) using the current interruption

method to correct for solution-phase IR drop external to the gas-fed electrode.

RESULTS

Using the alkaline solution technique, a sample of Pb2Ru206. 5 was prepared with

a reasonably high area (-35 m 2 g-i ) and some evidence of crystallinity (X-ray

diffraction) even without further heat treatment. After heat treatment at -350°C,

the X-ray diffraction becamevery well resolved. This result is important because

heretofore it has been assumed that in order to achieve high surface areas, it is

necessary to prepare the lead ruthenate with some substitution of the Ru with Pb (7-

8). Such "lead-rich" compounds are known to exhibit decreasing stability as the

amount of substitution increases (I). The 02 generation behavior of this

Pb2Ru206. 5 sample was significantly better than for other lower area samples of the

same stoichiometry that were examined with gas-fed electrodes.

Further characterization studies for Pb2Ru206. 5 are in progress in which the

effect of the heat treatment is being examined. The thermogravimetric analysis

(TGA), for example, indicates that -0.8 moles of water are present before any heat

treatment and are mostly removed at 800°C. The X-ray photoelectron spectra

(XPS), however, did not show discernable changes in peak shapes or binding energies

for Ru, Pb and O over the temperature range IO0°C to 500°C, in which most of the

water is lost.

139
Cyclic Voltammetry

The cyclic voltammetry (CV) for the Pb2Ru206.5 thin porous coating is very

complex. There is a large, gradually increasing anodic current on which are superimposed several small peaks (Fig. i, a-e). On the back sweep there is a large

cathodic peak. The charge under this peak (2.7 x 10 .2 C) corresponds to -12 _ of

the total possible charge expected for a single electron transfer to each Ru atom in

the coating (-0.8 mg or i.i x 10 .6 mol, yielding 0.214 C). From the crystal structure (9), there are 4 Ru atoms exposed on the face of each unit cell, whose edge is

10.253 A. This corresponds to -6.3 x I0 -I0 mol cm "2. -rThe BET area of the

Pb2Ru206. 5 sample used for the CV experiment was 6.0 m2 g-1 Therefore, for a 0.8

mg sample, the surface area was - 48 cm 2, corresponding to 3.02 x 10 .8 mol or 2.9 x

10 .3 c for a l-electron process for surface Ru atoms only. This value is one order

of magnitude less than the observed charge of 2.7 x 10 .2 C. One possible explanation is that -i0 monolayers are involved in a l-electron process. On the other

hand, for the valence states (pb2+)2(Ru4,5+)2(02-)6. 5, the following oxidation

processes are possible over the potential range:

4 e2 Pb 2+ "_ 2 Pb 4+ f

3 e2 Ru 4, 5+ ) 2 Ru 6+

for a total of 7 electron transferred. This would still account for -1.5 monolayer,

if all 7 electrons are involved. Therefore, it appears that the overall process

involves not only the surface but also the bulk (1.5 < x < i0 monolayers) of the

material. If the process is indeed a l-electron process, it may be assigned to the

Ru(IV)/Ru(V) transition proposed by Edgell et al (2). The existence of the smaller

peaks has not been explained thus far but may involve Ru atoms on the surface of the

oxide undergoing successive electron transfers with both Pb and Ru going to higher

valence states.

Oxygen reduction

The polarization curves for 02 reduction on a thin porous coating of Pb2Ru206. 5

at a series of rotation rates are shown in Fig. 2. The corresponding ring currents

for peroxide oxidation are very small, reaching a maximum of only -5 % of the disk

current (corrected for N) at the most negative potentials

Kinetic analysis of the ring-disk results using the method of Wr0blowa et al.

(i0) would lead to the conclusion that the 02 reduction is proceeding via

the direct 4-electron pathway to OH'. In addition, the slopes of the i "I vs.

f-i/2 plots (where f = rotation rate) correspond to an overall number of electrons being transferred of -3.8. The porosity of the electrode causes complications

in the interpretation of the data however. For exampie, if a catalyst is a go6_ i]

peroxide decomposer, it can appear that the 02 reduction is proceeding by a direct

4-electron pathway even though the reaction is proceeding by a series pathway, with

02 reduction to peroxide as the first step, followed by peroxide decomposition.

Overall this would yield 4 electrons. The reason for the possible confusion is that

the peroxide requires a relatively long time to diffuse out of the porous layer and

during this time has a good opportunity to decompose. Even so, the significant

potential range over which there is negligible peroxide picked up at the ring

rather strongly indicates a 4-electron process.

140
Tafel plot of the mass-transport corrected currents for f - 2500 rpm in the

potential range of +0.06 to -0.08 V is shown in Fig. 3. A linear relationship is

obtained with a Tafel slope of -0.063 V (decade) °I This suggests a non-charge

transfer step as rate determining following a fast outer-sphere electron transfer.

Oxygen reduction studies were also carried out over the pH range 11.9 to 13.9

keeping the ionic strength constant. The currents for 02 reduction increased with

decreasing concentration of OH'. The half wave potential also became increasingly

positive. The slopes of the mass transport-corrected Tafel plots were also almost

identical over the above pH range. A reaction order of -0.5 with respect to OHconcentration over this pH range was observed. This result together with the Tafel

slope of -0.06 V (decade) "I is consistent with the following mechanism for 02

reduction on the Pb2Ru206. 5 pyrochlore.

Ru +- (Ru 3+

Ru 4+_ _'_OH- + e- -------> \Ru 4+

+ OH- (Fast)

u 3+ (R u4+ -- 02- "+ 02 _ (Slow)

\Ru4+ Ru 4+

Ru4+ -- 02-

Ru4+

+ e- --------->

Ru4+ -- 02-

Ru 3+

__---->

Ru 4+ -- Ou4+ -- Ofollowed by a series of further steps resulting in the overall process

Ru + -- O" (Ru4_, OH -

+ 2 e" + 3H20 ------->

Ru 4+ -- O- \Ru4+1"

+ 2OHwhich would not generate solution-phase peroxide.

Ox_eneration

02 generation studies were carried out over a range of pH (ii to 14) at a

constant ionic strength. The catalyst was incorporated in a porous electrode

(geometric area = 4.8 cm 2) using Teflon T30B as the binder with a Ni screen current

collector and heat treated at 280°C for 2 h in an inert atmosphere. The Tafel slope

at room temperature varied from 0.032 to 0.038 V (decade) "I A reaction order of

1.0 was found with respect to OH" concentration over this pH range. A reaction

mechanism consistent with the reaction order is proposed as follows.

141
k s

Sz + OH- " (SOH) z+ + e" (Fast)

kb

(SOH) z+ ------> (SOH) z+l + e" (Slow)

(SOH) z+l OH- Sz+l -- 0 -- H o-- OH-

_(SOH) z+l OH- z+l __ _ __ H --- OHz+l OI -- H --- OH- _ <SZ+ + i + HOH

z+l _ -- H --- OH- Sz+ HOH

This mechanism is similar to the one proposed for 02 generation on RuO 2 electrodes

(Ii).

Gas-fed electrode measurements

The electrochemical behavior, including the 02 reduction and generation

activity as well as the stability in the anodic mode can be modified by substituting

a part of either the Pb or Ru with other metals. One such type of substitution

which was examined is that of Ir for Ru in the B site. Ir is expected to be more

resistent to anodic dissolution than Ru (13). The ionic radii are very similar (II)

and the lattice parameters for Pb2Ru206. 5 and Pb2Ir206. 5 are very similar (9).

Thus there is probably a continuous range of solid solutions possible. Horowitz et

al. mentioned that such compounds are possible to prepare using the alkaline

solution technique (7) but have not presented electrochemical results for compounds

of this type. Two compounds were prepared in our laboratory and the X-ray

diffraction indicates a single pyrochlore phase.

Although the effects on the 02 reduction are slight, they are encouraging

(Fig. 4). The effects on the 02 generation were also slight ( Fig. 5). It is not

clear at present to what extent these effects might be due to changes in the wetted

catalyst surface area.

Another approach to inhibiting the anodic dissolution of Ru is to use a

conductive solid ionomer either as replacement for the liquid electroIyte Within_th_

porous 02 cathode or as an overlayer on the solution side of the electrode. Even

though the dissolved species are anionic, (RuO42-, HPbO 2", PbO32- ) they should

diffuse relatively slowly in such ionomers due to size effects. Two types of

ionomers were used in this preliminary work. The first was a partially fluorinated

anion exchange membrane (RAI 4035)-as an overlayer. The second was a hydrogel

coating which consisted of a mixture of poly(dimethyldiallylammonium)chloride

(DMDAAC) and Nafion (14). The RAI membrane exhibited a definite effect in slowing

down the release of ruthenate into the solution when an electrode made from

Pb2[Rul.67Pb0.33]O6. 5 was in the 02 generation mode.

The pre-cast ionomer layer was not expected to have a significant effect on

142
the 02 reduction and this turned out to be the case (Fig. 6). With the hydrogel

coating, however, the ionomers can come into more intimate contact with the catalyst

in the porous layer and possibly improve the performance through an increase in the

02 concentration associated with the polymer fluorocarbon backbone. A small

improvement in the performance for 02 reduction at high current densities was

observed for the hydrogel-coated electrode and this is quite encouraging (Fig. 6).

With both ionomers there was an improvement in the 02 generation performance (Fig.

7). This was not expected but can be explained by pointing out that the inhibition

of the anodic dissolution would slow down the progressive changes in the surface

composition of the material that would occur during the dissolution.

ACKNOWLEDGEMENT

This research has been supported by NASA through the Lewis Laboratory and the

Department of Energy through a subcontract with Lawrence Berkeley Laboratory.

Samples of two pyrochlores were kindly supplied by Dr. H. S. Horowitz of Exxon.

The assistance of Mr. Wesley Aldred is acknowledged in carrying out the gas-fed

electrode measurements. Helpful discussions are also acknowledged with Dr. Martin

Shingler concerning the pyrochlore crystal structure.

REFERENCES

I .

2.

3.

,

5.

.

.

8.

9.

i0.

ii.

12.

13.

14.

H. S. Horowitz, J. M. Longo and H. H. Horowitz, J. Electrochem. Soc., 130, 1851

(1983).

R. G. Edgell, J. B. Goodenough, A. Hammett and C. C. Naish, J. Chem. Soc.,

Faraday Trans. I, 79, 893 (1983).

H. S. Horowitz, J. M. Longo, H. H. Horowitz and J. T. Lewandowski, in "Solid

State Chemistry in Catalysis", ACS symposium Series No. 279, R. K.

Grasselli and J. F. Brazdil, Editors, pp. 143-163, American Chemical

Society, Washington, D.C. (1985).

C. Iwakura, T. Edamoto and H. Tamura, Bull. Chem. Soc. Jpn., 59, 145 (1986).

J. A. R. Van Veen, J. M. Van Der Eijk, R. De Ruiter and S. Huizinga,

Electrochim. Acta, 33, 51 (1988).

J. Prakash, R.E. Carbonio, M. Razaq, D. Tryk and E. Yeager, 173rd Meeting of

the Electrochemical Society, Atlanta, GA, May, 1988, Ext. Abst., 88-1, 503

(1988).

H. S. Horowitz, J. M. Longo and J. T. Lewandowski, U. S. Patent 4,225,469,

September 30, 1980.

H. S. Horowitz, J. M. Longo and J. T. Lewandowski, Mat. Res. Bull., 16, 489

(1981).

J. M. Longo, P. M. Raccah and J. B. Goodenough, Mat. Res. Bull., 4, 191

(1969).

H. S. Wroblowa, Y. C. Pan and G. Razumney, J. Electroanal. Chem., 69, 195

(1976).

W. O'Grady, C. lwakura, J. Huang and E. Yeager, in "Proceedings of the

Symposium on Electrocatalysis", M. Breiter, Editor, pp. 286-302, The

Electrochemical Society, Pennington, NJ (1974).

R. D. Shannon and C. T. Prewitt, Acta Cryst., B25, 925 (1969).

M. Pourbaix, " Atlas of Electrochemical Equilibria in Aqueous Solutions," pp.

343-349, 373-377, Pergamon, Oxford (1966).

M. S. Hossain, D. Tryk, E. Yeager and A. Gordon, 171st Meeting of the

Electrochemical Society, Philadelphia, PA, May 1987, Ext. Abst., 87-1, pp. 466.

143
II

B

II

II

I

B

0

II.

II

.o-i

o

v

.'! ! § gs

II

o 0 o

i i

ORIGINAL PAGE IS

OF POOR QUALITY

is 111"

ill _ ii t,lit

: # # _!

• °

(Vii) JN_n9

g

o

o

r-1

o

!

o

r_

t

0

o

0

i

0

0

v_u'.LN31:II:InO

O

!

,ii

I

I

144

m.

o o

I"

I"

,6

>

0

>

<

Z

It) W

o;oO.
2.00 -

I,---4

I

,-I

#,---I

1.00

,:a

V

0.00 0

0

,...q

TAFEL SLOPE= -63 mV/DECADE

O o°

-0.05 -0.15 -0.25

E, VOLTS VS Hg/HgO

Fig. 3. Tafel plot for 02 reduction on Pb2Ru206.5 (thin porous coating) in O 2-

saturated 1 M KOH. The current values (in mA) for 2500 rpm from Fig. 2 were corrected for mass transport using I L = 1.75 mA.

I

O

d

>

-0.30 -

-0.20 -

-0.10 -

0.00 -

0.10

O SB / Pb2Ru206. 5

O SB / Pb2Rul.2Ir0.806.5

0.20 , , , ,

-2.0 -1.0 0.0 1.0 2.0

2

log i. m A/cm

3.0

Fig. 4. Polarization curves for 02 reduction with porous O2-fed (i atm) electrodes

in 5.5 M KOH at 25°C. The electrode contained 15.8 mg cm -2 pyrochlore, 14.6 mg cm -2

air-oxidized Shawinigan black and 12.2 mg cm -2 Teflon T30-B and was heat-treated at

280°C for 2 h in flowing helium.

145
Fig. 5.

I

O

d

t_

t_

>

>

0.60

0.50

0.40 -

r.d

-1.0

/_ SB / Pb2Ru206. 5

SB / Pb2Rul.2Iro.806. 5

O SB / Pb2Rul.sIro.506. 5

0.0 1.0

2

log i, mA/crn

2.0

Polarization curves for 02 generation with the porous electrodes of Fig. 4.

I

C)

t_

_0

>

>

-0.30

-0.20 -

-0.10

0.00

0.I0 -

SB / Pb2[Rul.67Pbo.33]O6.5

(Hydrogel)

0.20 , , , '

-2.0 -1.0 0.0 1.0 2.0

2

log i, mA/cm

3.0

Fig. 6. PolarizatiQn curves for 02 reduction with porous O2-fed (I arm) electrodes

in 5.5 M KOH at 25°C. The preparation of the hydrogel coating for curve 2 is

described in the te×t. The RAI membrane was pressed onto the solution side of the

electrode.

146
l

n_

O

_D

n_

o_

n2

>

>

0.60 -

0.50 -

SB / Pb2[Rul.67Pb0.33]06.5

_]No Membrane

O DMDAAC + Nafion

(Hydrogel)

_ RAI (4035_

r_; 0.40 " , ,

-1.0 0.0 1.0

2

log i, mA/cm

2.0

Fig. 7. Polarization curves for 0 2 generation with the porous electrodes of

Fig. 6.

147
=
N90-20469

OXYGEN ELECTRODES FOR RECHARGEABLE ALKALINE FUEL CELLS -

L. Swette and N. Kackley

Giner, Inc.

Waltham, Massachusetts 02254

I1"

The primary objective of this program is the investigation and development

of electrocatalysts and supports for the positive electrode of moderate temperature single-unit rechargeable alkaline fuel cells. Approximately six support

materials and five catalyst materials have been identified to date for further

development.

INTRODUCTION

Viable candidate materials for moderate temperature single-unit rechargeable alkaline fuel cells must meet the following requirements: i] good electrical conductivity (a more demanding requirement for supports than electrocatalysts); 2] high resistance to chemical corrosion and electrochemical oxidation

and/or reduction; 3] electrocatalysts, in addition, must exhibit high bifunctional electrocatalytie activity (02 evolution and reduction). Advanced development requires that the materials be prepared in high surface area forms, and

may also entail integration of various candidate materials, e.g., one or two

electrocatalysts distributed on a less active support material.

Candidate support materials have been drawn from transition metal carbides,

borides, nitrides (Ti, Zr, Hf, Nb) and oxides (La, Sr, Cr, Mo, W, Mn, Ni) which

have high conductivity (>i ohm-cm'l). Candidate catalyst materials have been

selected largely from metal oxides of the form ABO X (where A - Pb, Cd, Mn, Ti,

Zr, La, Sr, Na, and B = Pt, Pd, Ir, Ru, Ni, Co) which typically have been

investigated and/or developed for one function only, either 02 reduction or 02

evolution. The electrical conductivity requirement for catalysts may be lower,

especially if integrated with a higher conductivity support. For initial evaluation, materials have been purchased when available; subsequently, in-house

preparations have been attempted, to affect surface area and composition, if

necessary.

Candidate materials of acceptable conductivity are typically subjected to

corrosion testing in three steps. Preliminary corrosion testing consists of

exposure to 309 KOH at 80°C under oxygen for about 5 days. Materials that

survive chemical testing are examined for electrochemical corrosion activity;

the material is held at 1.4 V versus RHE in 30_ KOH at 80°C for 15 to 20 hours.

*This work is being supported by Nasa Lewis Research Center under the

direction of Dr. William Fielder, Contract No. NAS3-24635.

PRECEDING PAGE BLANK NOT FILMED

149
An acceptable anodic current is on the order of a few microamps/mg of material.

For more stringent corrosion testing, and for further evaluation of electrocatalysts (which generally show significant 02 evolution at 1.4 V), samples are held

at 1.6 V or 0.6 V for about i00 hours. The surviving materials are then physically and chemically analyzed for signs of degradation (visual examination,

electron microscopy, X-ray diffraction).

To evaluate the bifunctional oxygen activity of candidate catalysts,

Teflon-bonded electrodes are fabricated and tested in a floating electrode

configuration (Ref. I). Many of the experimental materials being studied have

required development of a customized electrode fabrication procedure. For

preliminary testing, catalysts of interest should show <500 mV polarization

(from 1.2 V) in either mode at 200 mA/cm 2. In advanced development, the goal is

to reduce the polarization to about 300-350 mV.

MATERIALS INVESTIGATED

Candidate Supports:

LaNiO 3 LiNiO x NbO 2

Candidate Electrocatalysts:

PbPdO 2 Pb2(Ir2.xPbx)O7_y

La.5Sr.5MnO 3 YBa2Cu30 x

Reference materials:

MoO 2 WO 2 HfB HfN TiN ZrN ZrC

Pb2(Ru2-xPbx)O7_y NaxPt304 La.5Sr 5Co03

PtTi230 x RuTil20 x RuMnl00 x RuMn20 x

02 reduction

10% Pt/Au (Johnson-Matthey, II m2/g)

10% Pt/Vulcan XC-72 Carbon (Johnson-Matthey, Pt: 130 m2/g)

02 evolution

Pt black (Englehard, 25 m2/g)

SOURCES OF CANDIDATE MATERIALS

Candidate materials for both catalysts and supports were purchased, if commercially available in powder form, as the most efficient approach for preliminary evaluation (electrical conductivity, chemical and electrochemical stability). Such materials offer the advantage of an economical purchase of_a material

of known purity in a quantity (5-25=g)_sufficient for preliminary qualification,

and particularly for disqualification. The disadvantage is that most of the

commercially avaiiable=materials have surface areas in the range from low

(<i0 m2/g) to very low (<i m2/g). Consequently, most commercial materials that

survive preliminary screening must be prepared in a higher surface area form,

in-house or by a custom fabricator, for effective evaluation as potential

catalysts and/or supports.

In many cases, especially for candidate catalysts, commercial materials

were not available. Preparation methods described in the literature, either

specific for the material or as a general model, were used when deemed appropriate to the material requirements in terms of electrical conductivity, surface

area, etc.

]50

Z
CHARACTERIZATION OF CANDIDATE SUPPORT MATERIALS

Candidate material preparations are typically analyzed by X-ray diffraction

(XRD) for chemical characterization. The objective of preparing fine powder

materials, however, is generally in conflict with obtaining sharp XRD patterns

(because of the line-broadening characteristic of high surface area powders);

thus the quality of these results is sometimes compromised. Firing materials

for a longer time or at a higher temperature usually increases the crystallinity

and improves XRD results, but at the expense of decreased surface area and some

uncertainty about the composition of the higher surface area material. In some

instances, materials of interest have been re-analyzed after extended corrosion

testing by XRD and/or scanning electron microscopy (SEM), to check for changes

in composition, reaction products and changes in morphology. SEM has also been

used occasionally to observe the particle size range of powders.

The electrical conductivity of candidate materials is estimated by compressing a small volume (e.g., 0.5-1 cc) of the powder at about 12,000 psi between

metal pistons within an insulating cylinder; the resistance of the powder is

measured directly, across the metal pistons, with an ohmmeter. If the resistivity thus measured is low (<5 ohm-cm), the resistance is redetermined by measuring the voltage drop across the powder under the flow of sufficient current to

generate easily measured current and voltage signals. The more accurate 4-point

method of measuring resistance has not been used because of the larger sample

volume requirement.

Surface areas of candidate materials are determined by the BET nitrogen

adsorption method using a Micromeritics Flowsorb II 2300 instrument. The samples are typically'outgassed at 150°C in the vent stream of the instrument, or

at -lO0°C in vacuum. All measurements reported have been made with 30% N 2 in He.

A summary of the measured physical characteristics of candidate materials

is presented in Table I. As a result of these measurements, the following

materials were eliminated from further consideration due to inadequate electrical conductivity: Nb02, PtTi23Ox, RuTil2Ox, RuMnl0Ox, and possibly La.sSr.sMnO 3.

The La 5Sr.5CoO 3 sample had a very low surface area but good electrical conductivity and thus remained a candidate for further testing. The LiNiO x sample had

both low surface area and marginal conductivity; XRD analysis, however, indicated a second phase of lithium carbonate (which might be leached out), consequently, this material was also retained for further testing.

STABILITY TESTING

A preliminary assessment of the chemical stability of the candidate support

materials is made by exposing the as-prepared powder to 30% KOH at 80°C under an

oxygen atmosphere. Indications of chemical reaction between the support material

and the KOH are color change of the solution or powder, evolution of gas, dissolution of the powder, weight change, and combinations of these phenomena. After

exposure for up to 5 days, the remaining powder is filtered out and weighed to

determine weight change, if any; small weight changes (e.g., <5%) have an equal

probability of being within experimental error for the measurement (weighing,

filtering, washing, drying, weighing), and need to be evaluated in conjunction

with other observations.

151
For an initial assessment of the electrochemical stability of candidate

support materials and catalysts, the steady-state anodic current is measured in

the range of 1.0 to 1.4 V versus RHE in 30% KOH at 80°C. The powder to be

tested is blended with polytetrafluoroethylene (PTFE, DuPont type 30 Teflon TM

suspension) at about 10-20% by weight and heated to 275-325°C, to try to achieve

a suitable compromise between physical integrity and good electrolyte penetration. A pure gold mesh is used as the current collector and the electrode is

suspended vertically in solution to prevent gas bubble occlusion of the surface.

If the anodie current observed, after initiation of potentiostatic control,

drops to the microamp range, the system is allowed to equilibrate overnight; the

steady-state anodic current is then recorded. In a second stage of testing

(-i00 hours), candidate materials are subjected to higher potentials (1.6 V)

representative of oxygen ev01utionconditions, and lower potentials (0.6 V)

representative of oxygen reduction conditions. The latter is intended to place

electrochemical stress on the materials used in the oxide form.

The value of residual anodic current measured by these methods is not an

unequivocal indicator of electrochemical stability. A low value of anodic

current (e.g., a few microamps/mg) is necessary but not sufficient to demonstrate corrosion resistance since the powder Nay passivate or deiaminate from

the current collector and exhibit a deceptive value. At the other extreme, a =,

high current may represent the onset of oxygen evolution rather than corrosion,

especially in the case of catalytic materials. Finally, the gold current collector always exhibits a base level of anodic current (5-i0 microamps/cm2).

Consequently, the anodic current values measured must be combined with other

observations such as weight loss or gain, color changes, microscopic examination

(SEM, TEM) and analysis (EDAX, _D, etc.).

Observations on chemical stability are summarized in Table II.

Electrochemical stability measurements are recorded in Tables III and IV.

Three candidate materials were eliminated from consideration due to evidence of=c_emicai _nstability-when the Sampie powders were exposed to KOH at

80°C under 02: HfB, HfN, and YBa2Cu30 x. Two additional materials were eliminated due to dissolution during electrochemical stress testing at 1.6 V: MoO 2 and

WO 2. As a result of these screening tests, the following materials remain as

potential candidates for further investigation, and may be classified as

follows:

Supports:

TiN, ZrC, ZrN

Catalytic Supports:

Electrocatalysts:

LaNiO 3, LiNiOx, La.5Sr.sCoO 3

PbPdO 2, Na.sPt304, Pb2(Irl.33Pb 67)O7_y,

Pb2(Rul.35Pb.65)O7.y, RuMn20 x

TM - Teflon is a trademark of E.I. DuPont Corp.

152
The following observations were made on these materials:

TiN: The coarse commercial powder _40 microns) and the low surface area

Giner, Inc. preparation (i micron, 2.3 mL/g) appeared to be quite stable up to

1.4 V vs. RHE. The latter material, held at 1.6 V for more than i00 hours, had

a slight surface discoloration, but otherwise, showed no evidence of reaction

(e.g., no weight change). Post-test XRD analysis indicated a strong TiN

pattern (a - 4.240 angstroms), no Ti02, and a few lines possibly due to a

silicate, phosphate or carbonate.

The higher surface area material (TiNx, Univ. of Calif.) similarly appeared

relatively stable in the chemical test (very small amount of white floc visible,

quantitative weight change measurement not successful) and anodically up to

1.4 V. Post-test XRD analysis showed a strong TiN pattern and three additional

lines that might be attributed to a phase such as potassium titanate. When the

electrode was held at 1.6 V, however, the material was largely lost, apparently

by dissolution. A gold-catalyzed TiN electrode was also found to be very reactive. After electrochemical testing, the surface of the electrode had degraded

to a refractory non-conductive material. Post-test XRD analysis again showed a

second phase (much stronger) consistent with the first analysis (potassium

titanate?). Based on these results, it appears that the non-stoichiometrlc

TiN x (x = 0.72-0.86), although promising in terms of the high surface area

achieved, is unstable at anodic potentials. Based on other results, TiN remains

a candidate support material under investigation.

ZrC: The commercially available material was a coarse (40 microns), low

surface area powder with about 2% Hf impurity. During the initial period of the

chemical stability test, gassing was observed; this was attributed to dissolution of the Hf impurity (weight change measurements were not successful). In

three repetitions of electrochemical testing, the anodic currents measured at

1.4 V were very low, 0.2 to 0.7 microamps/mg, equivalent to the background

current measured on the gold current collector. In testing at 1.6 V vs. RHE

from 50 to i00 hours, there was a visible loss of material, as well as a measurable weight loss (e.g., -40%), but this could not be distinguished from physical

shedding during oxygen evolution, and there was no other evidence of reaction.

This material remains a potential candidate support under development with

emphasis on higher surface area preparations.

Zr__NN: The commercially available material, a coarse, low surface area

powder, also appeared to be fairly inert in all stability tests. A weight loss

of about 9% was recorded for a 14-day chemical exposure test. The anodic current at 1.4 V was at a background level (0.2 microamps/mg) and there were no

signs of degradation after a week of oxygen evolution (1.6 V in 30% KOH at

80°C). This material also remains under development primarily to produce high

surface area powder.

LaNiO3: The commercial material has proven to be quite stable in all

corrosion testing. Since LaNiO 3 is also catalytic (Ref. 2) oxygen evolution

currents were measured at 1.4 V vs. RHE. After extended testing at 1.6 V and

0.6 V, XRD analysis still indicated a strong, unchanged pattern for LaNiO 3.

Attempts to prepare materials with surface areas greater than ~5 m2/g have been

unsuccessful to date.

153
LiNiOx: XRD analysis of the commercial sample (labeled "LiNi02" ) indicates

a strong second phase of lithium carbonate. This may account for the weight

loss recorded in chemical testing. It is anticipated that a more well-defined

material will be stable, somewhat catalytic, and a potential support material.

Higher surface area in-house preparations will be investigated.

L__a.bS_K.bCO___Q3: This commercial compound was quite conductive, but very low

in surface area. It showed a 28% weight loss in the chemical test but did not

appear to be reactive (or very catalytic) at anodic potentials. Higher surface

area preparations will be attempted for further investigation as a support.

PbPdQ2: The classical preparations of this material, by high temperature

firing of the mixed oxides (Ref. i0) or hydroxide co-precipitates, have generally resulted in fairly non-conductive powders (Ref. II), limiting our ability to

evaluate the material. It remains of interest because it has shown catalytic

activity on a carbon support (Ref. ii). A recent preparation by a novel method

has yielded a quite conductive material with the correct _b pattern for PbPdO 2.

In the limited chemical test performed to date, the material showed no weight

change or other signs of reaction, and will be investigated further.

N___a.8P___t304,prepared at Giner, Inc. (Batch #I), had moderately high surface

area and good conductivity. After 17 days of corrosion testing, primarily at

1.6 V vs. RHE, there were no signs of degradation or changes in appearance (no

XRD data)A and the electrode showed a similar level of anodic polarization at

200 mA/cm z (-400 mV) as a freshly-prepared electrode (-385 mV). This material

continues to ShOW promise as a bifunctional oxygen electrode catalyst, as

discussed in the next section. -

Pb2(Irl_33Pbi67_7_y: In our initial' investigation of this material

(Ref. Ii) it was found to be unstable in 30% ..... KOH at 80°C (e.g., 23 microamps/mg

at 1.0 V vs. RHE); at 1.6 V, the material showed substantial physical changes.

More recently, this same material was refired at 500°C. This had the effects of

improving its crystallinity and decreasing its reactivity below I_.4 V. The test

electrode survived 24 hours at 1.6 V, but delaminated from the current collector

at 0.6 V. It is anticipated that this material will be prepared in different

compositions and by alternative methods for further investigation.

P__hb2(RUl.35P__hb.65)O7_y, prepared in several batches by a different method

(Ref. 8) than the Pb-Ir pyrochlore above, was obtained in a crystalline form at

low temperature. The resulting materials, however, were unstable under most

test conditions, as evidenced by yellow coloration of all test solu£ions. Heat

treating the 3rd_atCh of this compound at 400°C resulted in a substantial reduction in surface area (73 --> 30 mr/g) and S£abilized the comp0und in the

chemical test (no coloration of the solution). This material is under active

investigation.

RuMn2Ox: This material has been subjected to very limited investigation to

date; it shows evidence of instability similar t0 that observed for the Ru

pyrochlore discussed above. Efforts are now concentrated on stabilizing this

catalyst by thermal treatments and alternative preparations.

=

=

=

=

=

_=

.=

=

k

m

l=

]54
OXYGEN ELECTRODE PERFORMANCE TESTING

Oxygen electrode performance testing is implemented in a floating electrode

cell (Ref. I) in 30% KOH at 80°C using a 1 cm x 1 cm electrode sample. Pure

oxygen is fed to the cell through a water presaturator. To fabricate a test

electrode, a powder sample is blended with PTFE (DuPont type 30 Teflon suspension) in the range of 15-40% by weight, as an approximate function of surface

area. The Teflon catalyst blend is then applied to a gold-plated Ni mesh with a

porous Teflon backing (l-micron pores), dried and thermally processed at 330-

360°C. It is frequently necessary to try other Teflon-catalyst compositions and

fabrication procedures to achieve an adequate balance in the hydrophobic/hydrophilic properties of the electrode.

The testing sequence is usually an oxygen reduction polarization test followed by an oxygen evolution polarization test, applying small potential steps

sufficient to yield a few data points in each log-decade of current density from

1 to I000 mA/cm 2. The potentials are controlled, and compensated for iR loss,

with a Princeton Applied Research model 173 potentiostat. Collecting oxygen

evolution data entails frequent interruptions to remove trapped gas bubbles.

After an oxygen evolution test, electrodes are sometimes too flooded to rerun an

oxygen reduction test. Development of an electrode structure adequate for bifunctional operation with these candidate electrocatalysts is an independent

research task; some experience, gained in the development of Pt black bifunctional hydrogen electrodes, will be beneficial to this effort.

The best performance data for each of the catalysts, as well as data for

the reference materials (10% Pt/Au, 10% Pt/C and Pt black), are shown in

Figure i. The following observations were recorded:

RuMn20 x exhibited exceptional oxygen evolution performance, and moderate

oxygen reduction performance, in the limited testing performed to date.

Unfortunately, this batch of material was not stable, as evidenced by the reddish coloration of the test solutions in this test and in the earlier corrosion

tests. We are currently investigating methods of stabilizing this catalyst by

thermal treatment and/or doping.

Pb2(Irl.33Pb.67)_Q7_y (fired at 400°C) gave good performance at lower current densities for both oxygen reduction and evolution, but polarized rapidly

above I00 mA/cm 2, suggesting a poor electrode structure. This preparation was

also unstable as discussed under stability testing. The sample refired at 500°C

was corrosion tested but not performance tested. Further preparation and testing of this catalyst is planned.

Pb2(RUl.35Pb.65)O7_y , prepared by the low temperature process described

(Ref. 8), showed from good (Batch #i) to exceptional (Batch #3) activity for

oxygen reduction, as presented in Figure 2. All of these preparations were

unstable in KOH, however. Post-thermal treatment at 400°C (Batch #3) stabilized

the material but resulted in a loss in performance, coincident with the loss in

surface area. Oxygen evolution performance was generally very poor with the

exception of Batch #i, as shown in Figure i. Performance of all materials was

typically transient and non-reproducible. This catalyst is still under

investigation with emphasis on stability and oxygen evolution performance.

155
N__%a.8P_tt304 - Performance data for two separate preparations of this catalyst

are shown in Figure I. The data for Batch #2 is the more recent and represents

improvements in electrode structure as well. This material is the best candidate bifunctional oxygen electrode catalyst developed to date in terms of both

stability and oxygen reduction/evolution performance. The emphasis in further

development is on increasing the surface area and preparing larger quantities of

material. Deposition on a support will be attempted also, when a suitable,

stable, high surface area support material has been developed.

Au/TiNx: Gold was deposited on the high surface area TiNx, by a Giner,

Inc. proprietary process, to investigate the feasibility of catalyzing this:_

support. The surface area was increased from about 38 m2/g to 51 m2/g, indicl _ :

ting an average particle size of about 25 nm. Electron microscopic examination

of this material at NASA, by Dr. William Fielder, indicated the presence of

particles in this range as well as larger agglomerates of particles. The

material could not be tested electrochemically because of the instability of the

TiNx, as discussed above, but indicates the potential feasibility of catalyst

deposition on this novel support.

ACKNOWLEDGEMENTS

We would 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

I. Giner, J.; and Smith, S.: A Simple Method for Measuring the Polarization of

Hydrophobic Gas Diffusion Electrodes. Electrochem. Technol., vol. 5, 1967,

p. 59.

2. otagawa, T.; and Bockris, J.O'M.: Lanthanum Nickelate as Electrocatalyst:

Oxygen Evolution. J. Electrochem. Soc., vol. 129, 1982, p. 2391.

3. Zhang, H-M.; Teraoka, Y.; and Yamazoe, N.: Preparation of Perovskite-type

Oxides with Large Surface areas by Citrate Process. Chemistry Letters,

1987, p. 665 .

4. Baythoun, M.S.G.; and Sale, F.R.: J. Mater. Sci., vol. 12, 1982, p. 2757.

5. Schwartz, K.B.; Prewitt, C.T.; Shannon, R.D.; Corliss, L.M.; Hastings, J.M.;

and Chamberland, B.L.: Neutron Powder Diffraction Study of Two Sodium

Oxides: Nal_oPt304 and Na0.73Pt304 . Acta Cryst., vol. 38, 1982, pp. 363-

368.

6. Shannon, R.D.; Gier, T.E.; Carcia, P.F.; Bierstedt, P.E.; Flippen, R.B.; and

Vega, A.J.: Synthesis and Properties of Platinum Metal Oxides of the Type

MxPt304. Inorg. Chem., vol. 21, no. 9, 1982, pp. 3372-3381.

7. Horowitz, H.S.; Longo, J.M.; and Lewandowski, J.T.: Lead Ruthenium Oxide,

Pb2[Ru2_xPb_+]O6.5 . Inorganic Synthesis, vol. 22, John-Wiley & Sons, New

York, 1983, pp. 69-72.

156
8. Horowitz, H.S.; Longo, J.M.; and Lewandowski, J.T.: Method of Making Lead

and Bismuth Pyrochlore Compounds Using an Alkaline Medium and at Least One

Solid Reactant Source. U.S. Patent No. 4,225,269, Sept. 30, 1980.

9. Adams, R.; and Shriner, R.L.: Platinum Oxide as a Catalyst in the Reduction

of Organic Compounds. III. Preparation and Properties of the Oxide of

Platinum Obtained by the Fusion of Chloroplatinic Acid with Sodium Nitrate.

J. Am. Chem. Soc., vol. 45, 1923, p. 2171.

I0. Lazarev, V.B.; and Shaplygin, I.S.: Properties and Structure of PbPdO 2.

Zh. Neorg. Khim., vol. 24, 1979, pp. 885-889.

ii. Swette, L.; and Giner, J.: Oxygen Electrodes for Rechargeable Alkaline Fuel

Cells. J. of Power Sources, vol. 22, 1988, pp. 399-408.

157
TABLE I - PHYSICAL CHARACTERIZATION OF CANDIDATE MATERIALS

Materlal

PbPdO 2

Na.8Pt304

Pb2 (Irl. 33 pb. 67) O7-y

Pb2 (RUl. 35 pb. 65) O7-y

PtTi230 x

RUTil2Ox

RUMnl0Ox

RuMn20 x

La.5Sr.5CoO 3

Source [Reference] Surface

Area

(m2/g)

Giner, Inc. #4 4

Giner, Inc. #I [5,6] 21

Giner, Inc. #2 [5,6] 17

Giner, Inc. #i (400°C) [7] 24

Giner, Inc. #I [8]

Giner, Inc. #2 [8]

Giner, Inc. #3 [8]

Giner, Inc. #3 (400°C) [8]

Giner, Inc. #i [9]

Giner, Inc. #I [9]

Giner, Inc. #i [9]

Giner, Inc. #i [9]

CheMaterials (Basic Vol.)

CheMaterials

55

35

73

3O

79

73

38

62

Elect.

Conduct.

(ohm-cm-1)

0.74

5O

56

4O

41

3O

33

26

(poor)

(poor)

(poor)

2.9

La.5Sr.5Mn03

YBa2Cu30 x Duke University 0.04 1.3

CheMaterials

Giner, Inc. #i [2]

Giner, Inc. #2 [3]

CheMaterials

Alfa Prod. (Morton-Thiokol)

Alfa Prod.

Alfa Prod.

Aesar (Johnson-Matthey)

Aesar

Alfa Prod.

Aesar

Alfa Prod.

Giner, Inc.

Univ. of California #2

Univ. of California #3

Univ. of California #5

Giner, Inc./U. Calif. #3

LaNiO 3

LiNio x

NbO 2

MoO 2

WO 2

HfB

HfN

ZrN

0.1 105

9 0.14

5

4

6

5-10

46-109

(poor)

N/M

0.8 0.4

N/M (poor)

N/M 3.4

N/M 3.5

N/M 2.1

1.7

0.3

N/M

2.3

56

38

21

ZrC (2% Hf)

TiN

Au/TiN 51

91

125

390

440

0.34

N/M

14

0.37

158
TABLE II - RESULTS OF STABILITY TESTS OF SELECTED CANDIDATE MATERIALS.

Codes: NVR - No Visible Reaction

SR - Slight Reaction (C = Color, G = Gas)

D - Dissolved

N/M - Not Measured

Material Source [Reference] Weight observaChange tions

(%)

pbPdO 2 Giner, Inc. #4 +0.7 NVR

Na.8Pt304 Giner, Inc. #1 [5,6] N/M NVR

Giner, Inc. #2 [5,6] N/M NVR

Pb2(Irl.33Pb.67)O7_y Giner, Inc. #i (400°C) [7] N/M NVR

Pb2(RUl.35Pb.65)O7_y Giner, Inc. #i [8] N/M SR-C

Giner, Inc. #2 [8] N/M SR-C

Giner, Inc. #3 [8] -20 SR-C

Siner, Inc. #3 (400°C) [8] N/M NVR

RuMn20 x Giner, Inc. #i [9] -14 SR-C

La.5Sr.5CoO 3 CheMaterials -28 NVR

La.5Sr.5MnO 3 CheMaterials -15 NVR

YBa2Cu30 x Duke University N/M SR-C

LaNiO 3 CheMaterials N/M NVR

Giner, Inc. #i [2] -25 NVR

LiNiO x CheMaterials -29 NVR

MoO 2 Alfa N/M NVR

WO 2 Alfa N/M NVR

HfB Aesar -i00 D-C/G

HfN Aesar -i00 D-C/G

ZrN Alfa -9 NVR

ZrC (2% Hf) Aesar N/M SR-G

TiN Alfa N/M NVR

Giner, Inc. N/M NVR

Univ. of California #2 N/M SR-C

159
m

N r'_ (_1 _1' O_ 0 I

G_ • _ • • I

_-_ o ,_r ,_r I

qP

I ,,..I o I I ,-I I

UO • I I I

,,,-I • _ ,,-I I I q-_ l

U 0

m

Z

0

,.-I ,---I I I

tJJJo

oo_o

_o_

,eoell

00_00

II _ I0 III _ _ 5_0_

II . I III

II 0 I_ |II 0 _ 00_

I II _ I_ III _ M _O_m

I II I Ill • •

,, _ ,_ ,,, o _ dd_g_

I _ II _ _ _IO _ _ _o_

l . lI • _- ,I_ -- -- .....

I o I I o o _ I o o oo_o_

A_=_.

oo

® _ oo _

q_

® d dd d d "_ "_' _

H _I H HH H H _.I _-IH

_:I I I

I'- _'-

0 _ _

I_ ¢;

I_ ;.i i_ _ 0

@ _ u_ 0

I # _ _ _ {J

M _ _ :_ 0 _"

• ¢'_ _ _ "

-M-M-M

000

-M-M-M

._

0_0_

000

_oo,

X

0

:_ 0

A

:m _) z

160 =
H

M

0

0

0

W

W

W

0

H

0

0

I

H

E._

0 _ _-" _-_

pa_

•d _ _ ._ .

CJ m_ 0 0 .. _I {/I {/l

Q

,Q II: C C C -,-4 -,-4

0 _." ID_ -_ .d

•,_ ._ _

_1 U_ 0 0

de ae a_ 0 ae

,,;i. _-i I _ rl I ".-- _ I

0

_ko o_e o_o

co o i.c)

_o • _o . • ._ •

'0 _ _ -- 0_I

•,-_ _ 'Cl _'_ _ _ _ _ •

:> _ . _ • _) _ -,_ • 0-,_ ,-_

0 _0 _00 -_ • ,t_

o Iw _ 0 0 ,'-i _ •-_-_ 0 _ _ _'

g

O

@

N

@

_w

@

0

0

0 0 _ _ "_-_

_ _0 o_ CC

O0

_=_ =_ _ -_ .,_

0 O0 (If _ _ oorJ

• _ • O0

,4

*H

k

I I

O O

g

a, _ _ _ _ 0 0

• e_ t_ _ • _ _ 0

o

0

w_

U_

oko

Z O

161
o

o

o

o

o

÷

o

'0 ÷_

D O

o

o

O4

+O

D _

0 o

o

+D _

O

d

0 _

+ D

o

o

+

o

_o

o

4

0

÷

0

+

O+

I>

o

4

D.

0

D_

0

r':_

0 0 0 0 0

1- © LO xl C9

0

0

0

o

I_ oo

4, _0

_oOo

q

4Do r_

4' 0 rl 0

_0_ o

,_o _o

• 4::

i:1

6

"t

13

13

0

m8

@

o

o

4

o

I1',4

O

O

T--

I'1

O

E

o

<

E

)-

Z

III

KI

Z

W

rr

rr

0

0 A

A 0

00 U_-

@_ vO

0 0

_1 ,--_0 00_

® +0o1>

_ .-IN (.1

Cggo

O'-"

_1 0 0 ,.q _"

114

•_+0_ <1

m

4

0 0 0 0 0 0 0 "_

OI _ 0 01 O0 I'-- _0

d o d d

m

=

=

=

m

(pe_oe_oo El! 'A) _HIzl "SA -1VIINEt_LOd

162
Ib

(> 0

II, D

oli,

0

• d

P i3

od

0 d

q

4

4

,.oB

o

4

o 0

13

ID

o <3

I_q3

@

j_

P

@0

@

P

1 t= I I I

0 0 0 0 0 0

c_ _- 0 O_ 00

.- _ o d d

0

0

o

0

0 _-

d

(pe_.oeJJoo _1! 'A) BH'c::I "SA OVIIN31Od

E

t_

<(

E

>_-

pZ

I11

a

pZ

LU

0E

rr

D

0

.o

Uo

0 o

oO

4J .-

_M

o,1:1

_n

oH

N _o

0_

@

@

_ x

.I._0

_0

O_

@

rj x

_0

O_

@

u_

@o

0

-H

163

N90-20470

NON-NOBLE ELECTROCATALYSTS FOR ALKALINE FUEL CELLS

S. Sarangapani, P. Lessner, M. Manoukian and J. Giner

Giner, Inc.

Waltham, Massachusetts 02254

INTRODUCTION

Noble metals, when used as electrocatalysts for oxygen reduction in metalair batteries and fuel cells, suffer from the disadvantages of high cost,

susceptibility to poisoning and sintering. Carbons activated with macrocyclics

have attracted increasing attention as alternative electrocatalysts for oxygen

reduction (Yeager, 1984). Initial activity of these catalysts is good, but

performance declines rapidly. Pyrolyzing the macrocyclic on the carbon support

leads to enhanced stability and the catalyst retains good activity (Wiesener,

1986). The exact nature of the catalytic sites after pyrolysis is still under

investigation (Scherson, et al., 1983; McBreen, et al., 1987).

Despite this promising method of treating macrocyclic activated carbon

catalysts, the performance decay is still too large to be acceptable. Part of

the stability problem occurs because the catalytic layer on the carbon surface

is, at most, a few monolayers thick. The approach described here is designed to

develop bulk doped catalysts with similar structures to pyrolyzed macrocyclic

catalysts. The transition metal and coordinated ligands are dispersed

throughout the bulk of the conductive carbon skeleton.

Two approaches to realizing this concept are being pursued, both involving

the doping of a carbon precursor followed by high temperature pyrolysis to form

a M-N-C catalyst. In one approach, the precursor is a solid phase ion exchange

resin. Two resins have been selected for doping. One is IRC-50 (Figure la)

which is a weak cation exchange resin that contains no nitrogen in its backbone

or exchange groups. The transition metal @nd nitrogen are introduced by either

exchanging a transition-metal (such as Fe +_ from FeCI 3) and nitrogen (from

• At

NH4OH ) sequentially or exchanging a chelation compound (such as Fe(blpy) 3 ).

The other ion exchange resin chosen was Chelite-C (Figure ib) which contains a

nitrogen-containing iminodiacetic acid exchange group. Exchange of a simple

transition metal salt results in metal-nitrogen coordination before pyrolysis.

In the second approach, gas phase precursors are mixed in a reactor and

pyrolyzed to form the catalyst. Acetylene (C2H 2) is used as the carbon source,

ammonia (NH 3) as the nitrogen source, and an iron-containing organometallic as

the transition metal source.

PRECEDING PAGE BLANK NOT FILMED

165
EXPERIMENTAL METHODS

Chelation compounds for doping the IRC-50 resin have included Fe(bipy)3SO4,

Co(bipy3)Cl 2, Fe(phen)3S04, Co(NH3)6CI 2 , and Co(en)3Cl 3. Doping is carried out

by mixing the resin with a solution of the dopant and agitating the mixture at

80°C for 3 hours. The doped resin is dried and then placed in a tube furnace.

The resin is pyrolyzed under an NH 3 atmosphere at temperatures near 900°C.

The gas phase precursors are pyrolyzed in a vertical tube furnace at

temperatures near ll00°C. The C2H2, NH3, and iron-containing organometallic are

introduced at the top of the furnace. Product is collected in a water trap at

the bottom of the furnace or inside the tube.

Electrodes are prepared from the bulk doped carbons by mixing the catalyst

with PTFE and spreading onto a Ag-plated Ni screen.

Electrochemical tests are performed using the floating electrode technique

(Giner and Smith, 1967). The electrolyte is 30% KOH and the temperature is held

at 80°C. The electrodes are tested galvanostatically and the measurements are

corrected for external ohmic losses. Potentials are reported versus a dynamic

hydrogen electrode (DHE).

The catalysts are analyzed for their carbon, hydrogen, nitrogen, and metal

content. Some catalysts have also had BET surface areas determined using a

Micromeritics Flowsorb II surface area analyzer.

Mossbauer spectroscopy has been used in an attempt to identify the species

present in the iron-containing catalysts. This work is being done in

conjunction with Professor Daniel Scherson at Case Western Reserve University.

RESULTS AND DISCUSSION

A. Electrochemical Tests

Figure 2 compares the electrochemical performance of catalysts prepared from

IRC-50 doped from solutions of Fe(bipy)3SO4, Co(en)3Cl3, Co(bipy)3SO4,

Co(NH3)6CI , and FeCI3-NH4OH with CoTMPP adsorbed and pyrolyzed on Vulcan XC-72.

There is a large improvement in performance on changing the dopant solution from

FeCI3-NH4OH to Fe(bipy)3SO 4 with the same resin and pyrolysis conditions. The

Fe(bipy), Co_en), and Co(hex) systems show a performance that is 50-80 mV less

at I00 mA/cm z than the CoTMPP on Vulcan XC-72. The surface areas of Co(en),

Co(hex) and Fe(bipy), doped carbons are in the same range as CoTMPP doped

Vulcan XC-72. The differences in the kinetic region may be attributed to the

differences in the surface concentrations of the catalyst. The earlier dropoff

of performance may be due to either ohmic losses in the structure or hindrance

to oxygen transport. The Co(bipy) system shows somewhat poorer performance;

stoichiometric analysis of the Co(bipy) system showed that only 1% Co was incorporated in the final catalyst (compared to 10% for iron) and, therefore, doping

conditions need to be adjusted.

The next set of experiments examines two different approaches for the

synthesis of "Fe-N" complexes. Figure 3 compares the performance of the IRC-50

Fe(bipy) system with Chelite-C doped with FeCI 3. In the first system, the Fe is

166
already coordinated with nitrogen in the bypyridine complex. Introducing this

complex into the ion exchange resin, allows bulk-doping and ensures molecular

level interaction between the carbon and the "Fe-N" complex. In the second

system, the resin has a backbone that contains nitrogen and coordinates with the

iron. In addition to differences in the density of the exchange sites (which

determines the metal/nitrogen ratio) in these two systems, the molecular interactions are bound to be different which mayresult in differences in oxygen

reduction performance. The carbon prepared from the chelating resin shows

essentially the sameperformance in the kinetic region, indicating that the

nature of the Fe-N complexes maybe similar regardless of their method of preparation, as long as covalent type interactions are ensured as opposed to purely

ionic type interactions. At higher current densities, there are somedifferences in the performance of these systems, which maybe attributed to the nonoptimized electrode preparation. It is also quite possible that the stabilities

of these complexes mayvary widely. Preliminary examination of the catalysts

using Mossbauer spectroscopy indicates that the IRC-50 and Chelite-C systems

contain one species in commonand one species each that is unique to that

catalyst.

Figure 4 shows the results of performance measurementson catalysts from

gas phase precursors. The undopedcarbon prepared from gas pyrolysis shows very

poor performance. Addition of the Fe and N has a dramatic effect on performance. The surface area of the Fe/N doped carbon in this figure, is only

19 m2/g and the nitrogen/metal mole ratio is 0.46. Studies are in progress to

improve both these factors.

B. Characterization of Catalysts

The catalysts are routinely analyzed for carbon, hydrogen, nitrogen, and

metal content. Table I shows the stoichiometry calculated for some typical

catalysts. There is still some unknown that is not accounted for in the C, H,

N, M. Some of this unknown is oxygen which comes from the COO" groups in the

resin and exposure of catalyst to the atmosphere. The oxygen is present in the

form of metal oxides. However, a calculation of the maximum weight percentage

of oxygen present in this form shows that it is not enough to account for all of

the unknown. The sample of IRC/Fe(bipy) was also sent for Si analysis. Some Si

is present, but Si or SiO 2 does not account for all of the unknown. Emission

spectroscopy was conducted to determine if any other heavy metals are present.

Only minor amounts were present. The sample was sent for sulfur analysis which

showed that some sulfur was present. This sulfur is probably from the sulfate

group in the Fe(bipy)3SO 4. This is because some of the dopant is imbibed in the

pores of the IRC-50 instead of being ionically bound to the COO" groups.

BET surface areas of catalysts prepared from the doped ion-exchange resin

precursors are in the range of 100-300 m2/g. When the IRC-50 resin with no

dopant was pyrolyzed, the surface area was less than 5 m2/g. This is consistent

with work in progress which indicates that the presence of the metal dopant has

a large effect on the nature of the pyrolysis process.

167
CONCLUSIONS

The doping of solid phase precursors followed by pyrolysis or the copyrolysis of gas phase precursors has allowed us to produce catalysts with good

activity toward oxygen reduction. Efforts are currently underway to better

understand the reasons for the catalytic activity of the bulk doped catalysts

with a view toward further improving their activity.

ACKNOWLEDGEMENT

Thi§ work is Supported by the=United States Department of Energy under

contract No_DE_AC02f87ERS0500.M001, We would ii_e to thank Professor

Daniel Scherson for providing the sample of CoTMPP adsorbed on Vulcan XC-72 and

performing the Mossbauer spectroscopy on the catalyst samples.

REFERENCES

Giner, J. and S. Smith. 1967. A Simple Method for Measuring the Polarization of

Hydrophobic Gas Diffusion Electrodes. Electrochem. Technol. 5:59.

McBreen, J.,_ W.E. O'Grady, D.E. Sayers, C.Y. Yang, and K.I. Pandya. 1987. An

EXAFS Study of Pyrolyzed Metal Macrocyclic Electrocatalysts. In Electrodes,

Materials and Processes for Energy Conversion and Storage, S. Srinivasan, S.

Wagner, H. Wroblowa (Eds.). Electrochemical Society Proceedings. 87-12.

Scherson, D.A., S.L. Gupta, C. Fierro, E.B. Yeager, M.E. Kordesch, J. Eldridge,

R.W. Hoffman, and J. Blue. 1983. Cobalt Tetramethoxyphenyi Porphyrin -

Emission Mossbauer Spectroscopy and 02 Reduction Electrochemical Studies.

Electrochim. A_t_ 28:1205.

Wiesener, K.

Reduction.

Yeager, E.

29:1527.

1986. N4-Chelates as Electrocatalyst for Cathodic Oxygen

Eiectrochim. Acta. 31:1073.

1984. Electrocatalysts for 02 Reduction. Electrochim. Acta.

TABLE I

Catalyst

IRC 50/Fe(bipy)

IRC 50/Co (bipy)

IRC 50/Co(en)

IRC 50/Co(hex)

Chel ite-C/Fe

Gas Phase Carbon

Gas Phase/Ferrocene Derivative

C

70.70

85.86

88.42

96.28

79.6

99.00

78.34

H

0.75

0.19

0.31

0.29

0.22

0.5

0.34

Wt%

N

2.36

0.210

0.620

0.i0

0.190

1.38

M

9.60

1.18

10.8

2.63

13.5

12.00

Unk

16.59

12.56

0

0.7

6.4

0.5

7.90

Surface

Area

(m'/g)

310.5

87.3

150.4

205.3

191.1

19

168
CH 3

r

_CCH2CHCH 2

------CHCH 2

Figure Ia: Structure of IRC-50 Weak

Acid Ion Exchange Resin.

A

L_

T

C3

>

>

E

1100

o

1000 'u

X

900 '+

800

700

600

500

0 0 0

X

+

_3

C

¢

4..o

o

Ck

 CHCH2CHCH

% - __CHC H E__r,C.j

%, ,o

400 " "

1

Figure lb: Structure of Chelite-C.

• . • • .| . . • ! .|

0 O0 0 00000

n aa O - 0 00000

DL• - 0 n n Oo_ o

÷ • •AAA • AAAAA (_ //ii "

++ + ++ AA A

+++ •

+

+ +++

+

2 3 4 _

..... I I I I .... II .....

10 100 1000

Current Density mA/cm 2

Fibre 2 : Polarisation curves for 02 reduction with porous electrodes in ?H KOH at 80"C.

Curve II 4.8_ w/w Co-TI4PP on XC-?2 carbon.

Curve 2:IRC-50 + Fe (added as bipyridyl complex).

Curve 3:IRC-50 + Co (added ae the hoxamino chloride).

Curve _: ZRG-S0 _ Co (added as tris othylonodiamin@ oobalt(IZl) chloride}. Curve • IRC-50 Co (added as bipyridyl complex).

Curve 6:IRC-50 + Fa (added as the chloride,) + N (added as NH4OH).

169
1100 - • ..... ' ....... ' .....

W

I

r_

>

>

£

E

¢

o

n

A

IiI

T

O

>

Ev

C

0

cl

F4guro 4 t

1000

g00

8OO

7OO

600

50O

'÷

+

÷ ÷ ÷

+ + +

÷ ÷ ÷

+ + +

÷

+

÷

-t-

+

400 ........ ' .......... L .......

1 i 0 100 1000

Current Density mA/cm 2

Figure 3: Polarization Curves for O_ Reduction With Porous Electrodes tn 7_ KOH at 80°C.

Curve 1: Chelite-C + Fe (_dded as the chloride) Pyrolyzed at 900UC. o

Curve 2:IRC-50 +Fe (added as the bipyrtdyl cm_lex), Pyrolyzed at 900 C.

..... - x L, _ L

1100 [" -_---- • ............ , ......

/

1000 [o

A

900 n _ _

0 O _ _ _A_ A_A_ _

0 z_&Z_z _

800 o

0 AIA AZ_ _

0

0 1

700 2

D

[]

O

6OO

n

500 []

[]

400 .................. ' ..... '

1 ......... 10 ......... -_ 100 : 1000

Current Density mA/cm _

Poltr£mat£on ourvom for 0 2 r oduat4on vith porou8 o]Lootrodo8 in 7M KOH &t I0"C.

C_rve lz Rleatrode va8 nado v£th pyrolyzod

CgH 2 ,IF Fo (&dded as dimothylamin0othyZferrooone).

Curve 2: pyz'olymed C2H 2 (1000"C).

]70
v v

N90-20471

CORROSION TESTING OF CANDIDATES FOR THE ALKALINE

FUEL CELL CATHODE

Joseph Singer and William L. Fielder

NASA Lewis Research Center

Cleveland, Ohio 44135

Current/voltage data have been obtained for specially made corrosion electrodes of

some oxides and of gold materials for the purpose of developing a screening test

of catalysts and supports for use at the cathode of the alkaline fuel cell. The

data consist of measurements of current at fixed potentials and cyclic

voltammograms. These data will have to be correlated with longtime performance

data in order to fully evaluate this approach to corrosion screening.

INTRODUCTION

Preliminary screening of materials intended for oxygen reduction performance

electrodes for the alkaline fuel cell is desirable in order to minimize the number

of long term test electrodes that might otherwise be required. ,Moreover, a

section from a performance electrode would not be as satisfactory as a specially

made corrosion electrode for this purpose because the latter should be made much

more wettable and intended for submersion in the electrolyte.

EXPERIMENTAL

I. Haterials

NiCo204 spinel was synthesized by co-precipitation, in the manner of Tseung

and Botejue (1985). This material was characterized in Singer et al (1987); its

surface area is 70 m2/g. _ The pyrochlore Pb2Ru206 5 was synthesized after

the method of Horowitz (1978, 1983); its surface area is lOl m_-/g. Two gold

standards, or blanks, were employed. One was the gold wire screen used as the

substrate for the oxides; the other was a piece from an actual working electrode

as u_ed in Orbiter fuel cells, comprising 20 mg/cm _ of Au-lO% Pt of surface area

12 mL/g and about 30 wt % PTFE (Teflon).

2. Electrode Fabrication

Pure Au wire screen was the substrate in the oxide electrodes. The screen was 50

mesh, of 0.25 in. wire; approximately 0.5 cm2 surface area of wire was included

in the submerged part of the oxide electrodes whose loadings were about 25

mg/cm 2. The IFC Orbiter electrode yielded a BET surface area of 2000 cm2 in

the electrode.

171
3. Measurements

4 mm x lO mm rectangular cuts of the fabricated electrodes were microtorch spot

welded to the 0.020 inch gold wire lead with a _ mm x 2 mm piece of Au foil

enfolding the clear area of the electrode and the wire. The contacted electrode

was immersed in the floating half-cell of Giner (1967) within the cylindrical cage

into which the reference electrode faces (Fig. l). The immersion included all the

active material and a minimum of the clean gold above it. The Teflon bracket of

the half-cell shown in Figure l which is used to contact the dry side of working

electrodes, was completely removed. The electrolyte was 30% KOH, the counter

electrode was Pt or Au foil, and N2 with about l ppm 02 was passed through a

30% KOH bubbler before the cell. Most measurements were made at 240 C.

Stirring was found to be detrimental.

Currents at fixed voltages were measured with a PAR 173 potentiostat with a 376

attachment leading into a recorder or a voltammeter. The 02 reduction range

between l.O and O.6v (RHE) was of most interest for observing corrodibility,

although some data were taken beyond these limits.

RESULTS

The current obtained from the gold wire screen, which was used as the

collector-substrate for all the materials except the IFC electrode, is shown in

Figure 2. The rapid drop to a steady current of a few microamps is noted. The

log-log plot was found generally useful for extrapolation to long times.

The IFC Au-Pt electrode, based on a pure gold wire screen of denser mesh than used

in the other electrodes, required a longer time to fall to the same microamperage

than required by the bare gold wire (Figure 3). Figure 4 shows the more rapid

fall to a lower current for the same electrode at O.6v than at O.3v.

The cyclic voltammograms of Figures 5 and 6 show appreciable differences between

the bare gold wire and the Orbiter teflonated gold electrode.

Corrosion currents obtained with the Pb-Ru pyrochlore and the Ni-Co spinel

electrodes are shown, respectively, in Figures 7 and 8. Cyclic voltammograms for

these i_o specimens are given in Figures 9, lO, and If.

DISCUSSION

The final currents obtained with the two gold standards are of the order of

microamps. It may be assumed that most of this is due to trace 02 in the

environment. The IFC electrode may be expected to be most efficient at reacting

with trace 02, having been optimized as a porous hydrophobic-backed reduction

electrode. Perhaps another type of gold electrode standard could be envisaged,

namely a corrosion type designed for maximum wetting, using pure Au colloid and

minimum Tefl on.

The cyclic voltammogr_ms of the gold specimens show sensitivity to absorption of

trace oxygen species at 0.925v (RHE) = 0 volts (Hg/HgO); the IFC electrode plainly

shows, in addition, its Pt content, in its resemblance to Pt cyclic voltammograms

throughout the literature (Will, 1966). The sharp rise in current above 500v

(Hg/HgO) is due to 02 evolution, which itself would supply 02_ to both gold

172
electrodes in addition to the ca. 1 ppmin the N2 gas supply, and to back

diffusion and leakage of air. In summary,the very low currents rapidly obtained

in Figures 2, 3, and 4 indicate a satisfactory standard for relative

non-corrodibi I i ty.

The pyrochlore used in this work is part of a study to makean improved material

with the expectation of improving its stability reproducibly. Whether the present

data indicate progress in this direction cannot yet be stated; however, Figures 7,

9, and lO obtained with this pyrochlore synthesis showno drastic corrosion, in

agreementwith Horowitz (1983). This surmise is fortified by comparison with the

spinel data of Figures 9 and II. NiCo204 has been reported to be unstable to

reduction potentials (King and Tseung (1974) which is one reason it was used in

this study. During the performance represented by Fig. 8 a black coating

developed on the counter electrode. This was dissolved in HCI. Ni and Co were

found, with about 6 times as much Co as Ni; however, no attempt was made to relate

this finding quantitatively to the corrosion data.

Comparisons of corrosion tests, like these reported, with longtime performance

electrode tests, will have to be made to evaluate this type of screening. The

present preliminary work suggests that this approach is worth pursuing.

CONCLUSIONS

Corrosion test screening of candidates for the oxygen reduction electrode of the

alkaline fuel cell has been applied to two substances, the pyrochlore

Pb2Ru206. 5 and the spinel _iCo204. The substrate gold screen and a

sample of the IFC Orbiter Pt-Au performance electrode were included as blanks.

The pyrochlore data indicate relative stability, although nothing yet can be said

about long term stability. The spinel was plainly unstable. For this type of

testing to be validated, comparisons will have to be made with long term

performance tests.

Ack nowl edgement

The collaboration of PSI Technology Company of Andover, MA, under Contract

NAS3-25199, in the preparation of the oxides and their electrodes used in this

study is herewith acknowledged. The provision of a sample of Orbiter electrode by

Dr. Ron ;4artin of IFC is gratefully acknowledged.

REFERENCES

Giner, J. and Smith, S. (1967): Electrochem. Tech. 5_, 59.

Horowitz, H. S., Longo, J. M., and Lewandowski, J. T. (197S): U. S. Patent 4,

129, 5P.5.

Horowitz, H. S., Longo, J. M., and Horowitz, H. H. (1983): J. Electrochem. Soc.

130, 1851.

King, W. J. and Tseung, A. C. C. (1974): Electrochem. Acta, 493.

Singer, J. and Fielder, W. L. (1987): NASA TI_ I00239.

173
Tseung, A. C. C. and Botejue, J. (1986): Int. J. Hydrogen Energy l_]_l,125.

Will, F. (1966): J. Electrochem. Soc. ll?, 481.

]74
PRESATUILATOR_o__ X [

_E._ I 1

ELECTRO_ _. "--J

etF6,L21°o_-__i_ COUNTEAE

NG E

vU_,U.U.tJ _--"

VOAKINGELECTRODE

P! COUNTER

ELECTRODE

LEADWIRE "_ ELECTROLYTE

LEVEL

FOIL COUNT-tRELECTRODE

FIG. 1. FLOATING HALF-CELL

n.

e:

z

u

10 -1

10 -2

10 -3

10 -4

10 -5

10 -6

| I w

I I |

10 2 10 3 10 4

TIME, Sec

Gold Wire Screen 0.300 V(RHE); 24 deg C

FIG. 2. CORROSION CURRENT

175
lo-li

10 -2

. 10 -3

i i0_4

10 -5

10"6

IFC Au-Pt ORBITER ELECTRODE

I I I

I t I

102 103 104

TIME, Sec

8.300 V(RHE); 24 deg. C

FIG. 3. CORROSION CURRENT

io -I

lO -2

10_3

z_ 10-4

10_5

10 -6

IFC Au-Pt ORBITER ELECTRODE

I I I

,I i |

102 103 104

TiHE, sec

0.600 V(RHE); 24 deg. C

FIG. 4. CORROSION CURRENT

176
GOLD WIRE SCREEN

a.

=!

I

e=

N

vlN

w

200

100

f

t f f ( f ,, ¢ ! ..L r r f ¢

-500 0 500

E vs Hg/HgO (mv)

Sweep: 600 to -826 my, 1 my/s, 60 deg. C, Nitrogen, 30%KOH

FIG. 5

[

CYCLIC VOLTAMMOGRAM; GOLD WIRE SCREEN

IFC Au-Pt ORBITER ELECTRODE

=oo j

=.

'_ 100 f

f |_. | I E i | ! i | i ! | ,I

-500 0 600

E Ve HoIHgO (my)

Sweep- 51_4 1o -706 my, "Jmy/s, 24 deo.C, NllroOen, 30%KOH

FIG. 6 CYCLIC VOLTAMMOGRAM

177
10-1

10 -2

, 10 -3

i 10_4

10"5

i0"6

10 -1

10 -2

. 10 -3

Z

10-4

U

10 -5

10 -6

.PSI Pb-Ru PYROCHLORE CORROSION ELECTRODE

I ! I

I I I

10 2 10 3 10 4

TIME, Sec

0.700 V(RHE); 24 dbg. C

FIG. 7 CORROSION CURRENT

PSI Ni-Co CORROSION ELECTRODE

i i I

I w I

102 103 104

TIME, Sec

0.60OVIRHE) L 24 deg. C.

FIG. 8 CORROSION CURRENT

178
PSI Pb-Ru Pyrochlore

u)

n

:!

,I[

P)

I

o.

v

_o

v-

)4

i

20

10

-10

F .._..,.._,.,--.r

f

L : I_ !

-300 -2O0 -100 0 100

E vs HgtHgO (my)

Sweep: 74 to -326 my', 1 mvle, Nitrogen, 30% KOH

FIG. 9 CYCLIC VOLTAMMOGRAM

PSI Pb-Ru Pyrochlore

in

D.

=E

'( 0

f,)

!

a.

-50 M

50 .

_,_..r.,.___.r _ -'w'''P-_

/f f

! ! I I I . I , ,,J_ i .

-500" 0

E vs Hg/HgO (my)

Sweep: 74 to -726 my, 1 my;s, 24 deg. C, Nitrogen, 30% KOH

FIG. 10 CYCLIC VOLTAMMOGRAM

179
NI-Co SPINEL CORROSION ELECTRODE

10

n

=i

!

X

e

o b

"t

-10 !

-4OO --aO0 -200 -100 0 100

E vs Hg/HgO (mv)

Sweep: 74 to -326 mv,80 deg. C, Nitrogen, 30% KOH

FIG. 11 CYCLIC VOLTAMMOGRAM

]80
N90-20472

STRUCTURAL COMPARISON OF NICKEL ELECTRODES AND PRECURSOR PHASES*

Bahne C. Cornilsen, Xiaoyin Shan and Patricia Loyselle

Michigan Technological University

Houghton, Michigan 49931

In this paper we summarize our previous Raman spectroscopic

results and discuss important structural differences in the various

phases of active mass and active mass precursors. Raman spectra

provide unique signatures for these phases, and allow one to

distinguish each phase, even when the compound is amorphous to x-rays

(i.e. does not scatter x-rays because of a lack of order and/or small

particle size). The structural changes incurred during formation,

charge and discharge, cobalt addition, and aging will be discussed.

The oxidation states and dopant contents are explained in terms of the

nonstoichiometric structures.

INTRODUCTION

Nickel electrode active mass has traditionally been classified in

terms of two distinct cycles, defined in equations 1 and 2 (ref. i).

_-Ni (OH) 2 .... > 7-NiOOH (i)

laging T overcharging

_-Ni (OH) 2 .... > _-NiOOH (2 )

The four end-member phases have been named on the basis of x-ray

diffraction pattern similarities between electrochemical materials and

chemically prepared materials, _-Ni(OH) 2, _-Ni(OH) 2, and 7-NiOOH

(ref. 2).

Our research has shown that these materials can be distinguished

using Raman spectroscopic analysis (ref. 3,4). The unique spectral

signatures found for the various phases allow definition of these

structures and reflect significant variations. The two cycles in

equations 1 and 2 differ in degree of nickel deficit nonstoichiometry.

The nonstoichiometric structural model proposed in reference 3

provides a framework in which the electrochemical properties can be

understood. The lattice structures of active mass are non-close

* This work was supported by NASA Lewis Research Center

Grant No. NAG 3-519

181
packed. The structural difference between the a/7 and the #/_ cycles

as not previously understood (ref. 5). #-Ni(OH) 2 , the traditional

close packed model of discharged active mass, is-not present in active

mass. Reanalysis of the x-ray diffraction patterns for charged active

mass and 7-NiOOH has supported a non-close packed structure, nickel

oxyhydroxide type structure, in agreement with vibrational spectra

(ref. 4). We define structural details and inter-relations between

these materials, including the structures of the cobalt containing

cathodic-_ and active mass (ref. 3,6).

To distinguish the active mass structural phases from the other

structures, we shall use prefixes, 2_, 37, _#, and _# (ref. 3). This

allows retention of the historical name indicative of intra-layer

structure measured by the x-ray patterns, but also allows aeiineation

of important structural differences. The structures of non-close

packed active mass are defined in terms of NiO 2 layers which exhibit a

nickel deflcit, and therefore can contain foreign cations, and

interlamellar cation sites. These are defined in terms of atom sites

present in the R_m, NiOOH type lattice (ref, 4). The_

nonstoichiometric formulae written in this paper express these two

types of sites, nickel or proton sites, and the dopants 6n them, by

writing the former ahead of the oxygen pair and the latter after the

oxygen pair.

EXPERIMENTAL

Preparati0ns of the cathodic-_, ordered #-Ni(OH)2 , disordered _-

Ni(OH)2 , and active mass phases (2_, 37, 2#, or 3_) have been

descriSed previously (ref. 3,4). Chemical analyses indicate that the

disordered E-phases exhibit a nickel deficit (x < 0.16) with empirical

formulae consistent with Ni I . (2H)x(OH)2. The empirical formula for

cathodlc-_ was obtained in t-he same manner.

Cobalt (5%) containing cathodic-_ phase was electrochemically

deposited on nickel foil (i0 min.), f_om aqueous nitrate solution (0.5

M), at a current density of 4.6 mA/cm _. Cobalt containing active mass

has been produced by charging these thin films in 0.5 M KOH. The low

KOH concentration minimizes aging. Cyclic voltammetry was used to

cycle the films between 0.25 and 0.55 V (with respect to a

0.5 M Hg/HgO reference electrode) at a 7.8 mV/min, scan rate.

The spectra and structures for the above active mass are

comparable to those found in commercial electrodes cycied to 80% depth

of discharge (DOD) and charged to 120% of full charge (20% excess

accounts for oxygen evolution during charge). Active mass is defined

as the material or materials found in a "formed" electrode which are

reproducibly observed during charge and discharge cycling.

Raman spectra were collected with a Ramanor HG.2S

Spectrophotometer with photon counting electronics. Multi-scan

spectra were required to provide adequate signal-to-noise ratio. The

182
excitation was the 514.5 nm line from an argon ion laser (10-20 mW on

the sample).

RESULTS AND DISCUSSION

The band positions in the Raman spectra of active mass

precursors, in active mass itself, or other related phases are

summarized in Table I (ref. 3,4). These spectra provide unique

signatures allowing distinction of each phase. The basic structural

unit in each of these materials is the NiO 2 layer, which contains

interstitial, octahedral nickel atoms wlthln a pair of close packed

oxygen planes. For the ideal crystal lattice the proton sites are

interlamellar (between NiO 2 layers); however, for a defective crystal,

protons may occur within an NiO 2 layer (possibly on or near empty

nickel atom sites). The structural relationship between two NiO 2

layers (layer stacking), the proton positions, and the internal

structure for each NiO_ layer influence the space group symmetry and

therefore the spectral-selection rules. The spectral selection rules

control the number and relative intensitSes of the Raman and infrared

vibrational bands. The 4000 to 2000 cm -_ region of the spectrum

contains vibrational modes which reflect the symmetry and bonding of

protons to oxygen atoms. The 600 to 200 cm -x region of the spectrum

contains Ni-O lattice modes, involving motion within the twodimensional layers. The number of lattice modes can be influenced by

layer stacking interactions. The wavenumber positions of vibrational

modes depend on the chemical structure and bond strengths, and

therefore vary with oxidation state and/or nonstoichiometry.

Ordered and Disordered _-Ni(OH) 2

The _-Ni(OH) 2 structure is the best characterized, being the only

compound in this series whlch has been studied in single crystal form

(ref. 7,8). The Raman spectrum (Table I) of this highly ordered,

crystalline phase contains only one O-H stretch and three lattice

modes (one is very weak) (ref. 4). This spectrum agrees with

theoretical selection rules predicted on the basis of the known

crystal structure. The NiO 2 layer stacking is known to be close

packed in thls structure. The three lattice modes originate from the

intra-layer Ni-O modes and from inter-layer interaction between the

close packed layers present in this stacking sequence (ABAB).

A disordered form of _-Ni(OH) exhibits a Raman spectrum which is

quite distinct in comparison with _he above ordered compound (ref. 4).

It contains two additional O-H stretching modes and one additional

lattice mode (see Table I), but is otherwise identical. This shows

that the basic structure is that of _-Ni(OH) 2, but with additional

features which relate to an empirical formula with Ni/OH < 0.5. A

structural model of a hydrated E-phase with interlamellar water can be

written as shown in equation 3.

]83
Ni (OH) 2" YH20 (3)

The lack of molecular water vibrations in both the IR and Raman

spectra suggest this hydrated model is inadequate. The lattice

parameters observed for these disordered-_ materials are comparable to

those of ordered _-Ni(OH)2 , and do not provide room for interlamellar

H20 molecules. Therefore this model is inadequate.

An alternative formulation with an equivalent stoichiometry,

where x = y/(2 + y), can be written (equation 4).

Nil_ x (2H) x (OH) 2 (4)

Nickel vacancy disorder is proposed based_on the empirical nickel

deficit (x < 0.16) and the vibrational spectra. A second type of

disorder is possible. Layer stacking disorder is likely, based upon

x-ray line broadening (ref. 9). Recent EXAFS analyses support the

nonstoichiometric model (ref. i0).

Charged Active Mass and Discharged Active Mass

Thechemically oxidized 7-NiOOH structure is the prototype for

charged active mass (ref. 4). This non-close packed NiOOH type

structure contrasts sharply with the close packed #-Ni(OH)2 structure.

The two lattice modes (Table I) and the lack of an O-H stretching mode

originate from the non-close packed layer stacking (ABBCCA) and

interlamellar hydrogens on centric lattice sites. This structure has

been supported by reinterpretation of the x-ray diffraction pattern

(ref. 4).

Both forms of charged active mass (3_- and 37-NiOOH) exhibit the

same selection rules which indicates that they have similar structures

(ref. 3). This crystal structure together with the point defect

structure allow one to understand the origin of the variable formulae

and oxidation states. The key difference between the 3#- and 37-NiOOH

phases is the magnitude of the nickel deficit and £he resuitant point

defect structures (see Table II). The nickel deficit is constant

within each cycle. The 37-phase generally contains potassium as well.

The amount of potassium uptake is determined by the nickel deficit,

i.e. the nickel vacancies are filled with potassium. The

ramifications of this nonstoichiometric structure upon the

electrochemical properties are significant (ref. 3). The differing

oxidation states allow the _/7 cycle to have a greater capacity. The

nonstoichiometry and point defects also control proton diffusion and

el ectron conduct ion. _ __ _ _ = __

b_schargea active mess (2_" and 2_£phaSes) dis_iays_spec£ra and

selection rules similar to those for charged active mass (ref. 4). We

propose a s_ar non-close packed layer structure. The nickel

deficit in each is comparable to that in the respective charged phase

(see Table II).

184
The wavenumber positions vary with charge state and defect

content (see Table I) as the active mass is charged and discharged.

The a/7 bands are higher in energy (wavenumber position) than the 8/8

bands. This shift parallels the increased nickel deficit observed for

the a/7 phases. Therefore, the change in defect structure is also

reflected in the band positions. The energy of a vibrational mode is

proportional to the bond strength (i.e. force constant). An increase

in oxidation state normally leads to an increase in wavenumber (ref.

Ii). The mechanism by which the nickel deficit influences wavenumber

is the same, in that the change in nonstoichiometry induces an

increase in average oxidation state and thereby, an increase in

wavenumber.

The traditional structures proposed for discharged active mass

(2a and 28) are 8-Ni(OH) 2 or the so-called a-phase structure (ref. 2).

The Raman spectra of discharged active mass clearly demonstrate that

formed and cycled materials are not isostructural with 8-Ni(OH) 2

(compare spectra in Table I) (ref. 3). The Raman spectrum of a

chemically precipitated a-phase also differs from that of discharged

active mass (see Table I). Therefore, the structures of chemical-a

and 2a or 28 differ as well. The structural difference between

discharged active mass (2a) and a cathodic-a phase precursor is more

subtle and will be discussed below.

Implications of the Nonstoichiometric Structure of Active Mass

The oxidation state maximum seen for 37 is +3.67. This oxidation

state is predicted for a defect structure which incorporates potassium

on 0.25 vacant nickel sites. With a +i dopant (K) on the nominal

Ni(III) site, only 2 electron holes remain per site to ionize and to

form two Ni(IV) species. Ionization provides 0.5 moles of Ni(IV),

equation 5, giving an average oxidation state of 3.67 for the total

0.75 moles of nickel per mole of compound.

Ni +4 Ni +3 K + OOH (5) 0.5 0.25 ( )0.25

Discharge of this material reduces the nickel, but the 0.25 protons on

the nickel vacancies produce an average nickel oxidation state of 2.25

for the 2a-phase (see Table II). It is apparent that the nickel

cannot be reduced completely to 2.00 because the lattice requires

charge balance and the vacant nickel site (nominally 2+ on the

discharged lattice) holds only one proton (based upon the empirical

formula).

The 37 point defect structural model explains how Ni(IV) ions can

be introduced into the nominally Ni(III), NiOOH-type crystal lattice.

In order for the average oxidation state of nickel to exceed 3+, a

nickel deficit is required. The 38 structure (Table II) demonstrates

that the nickel deficit is not the only prerequisite to an oxidation

]85
state above 3+. The dopant content, three protons on the nickel

vacancy, reduces the average oxidation state to 2.9 in this case.

It has been observed that the nickel vacancies are invariably

filled with protons or alkali metals, and that the excess monovalent

cations fit exactly on the vacancies, within reasonable experimental

error. When alkali metals are low, one to three protons can be found

per vacancy. When sufficient alkali is present in the formula, it

generally fills the vacancy. Therefore, when writing formulae, we

have chosen to place potassium or sodium on the nickel vacancy sites

in preference to interlamellar sites, We do this to demonstrate

correlation of potassium content and vacan6y_content. This

observation appears to suggest that the potassium on nickel site

positioning is fact; however, it is important to realize that this may

be coincidence. We have no positive, experimental evidence to prove

this, nor to disprove it. It can be said that the potassium does

stabilize the nickel deficit.

Thus the point defect structure (nickel deficit and dopant K or H

contents) controls both the maximum oxidation state and the minimum

oxidation state. Needless to say, the oxidation state change between

the two extremes controls the electrode capacity. The same logic

applies to the _/_ cycle, with the result _ the _ narrower oxidation

state range leads to a lower capacity. Therefore, the

nonstoichiometric model allows one to understand the variable

oxidation states of nickel active mass as well as the variable

capacitieS.

Cathodic-_ _Precursors of Active Mass

The two Raman lattice m0_es observed for cathodically deposited

electrode-precursor materials are consistent with a non-close packed

structure. Although the cathodically deposited o'phase materials

exhibit a broad O-H stretching mode, the lattice mode region differs

little from that of discharged active mass (Table I) (ref. 4).

Therefore the two structures are related in terms of layer stacking.

X-ray powder diffraction data also reflects this similarity (ref.

i, 2). However, the O-H stretching mode indicates that there are

differences in wa_er content an_/0r proton b0nding. This structural

difference may be significant.

The lattice mode positions (see Table T) for the cathodic-_ phase

are lower than for the discharged active mass phases (2_ and 2_).

This shift is related to the difference in bond strengths associated

with a small change in oxidation state. The cathodic-_ material is

green and is Ni(II). The oxidation states of the dark discharged

active mass are greater than 2 (see Table III).

The empirical formula of a cathodic-_ material is written in Table II.

The nickel-oxygen ratio in these materials is generally lower than for

discharged active mass. This can be caused by either additional

=

]86
molecular water (equation 6) or a more extensive nickel deficit

(equation 7).

Ni0.75 (2H)0.25 (OH)2"0-33 H20

Ni0.65 (2H) 0.35 (0H) 2

We favor the molecular water explanation because the Raman spectrum of

cathodic-_ displays O-H stretching modes which are typical of water.

The hypothetical nickel deficit in equation 7 is also greater than

0.25; the 0.35 value is larger than seen in other materials to date.

The above results show that discharged active mass, charged

active mass, and cathodic-_ phases are non-close packed and related in

structure. The proposal of this structure for the _-phases disagrees

with literature which suggests close packed inter-layer stacking (ref.

i). The literature powder pattern has been assigned a low reliability

in the JCPDS file (ref. 12) and has been questioned in the literature

(ref. 13). Further research is needed to understand how the excess

"water" is incorporated structurally.

Cobalt-Containing Nickel Electrodes: Precursors and Active Mass

Raman spectra of cobalt containing precursors, i.e. cathodically

deposited materials, are distinct and differ from that of cobalt-free,

cathodic-_ phase (ref. 6). The Raman spectrum of a 5% cobalt,

cathodic-_ displays two Raman bands at 460 and 530 cm -I (Figure i).

The wavenumber positions are identical to those observed with no

cobalt, but the intensity ratio for the two band_ is inverted. A very

weak O-H stretching mode is observed at 3652 cm- . The weakness of

this band and the intensity reversal distinguish the cobalt-free and

cobalt containing spectra. These changes reflect the significant

influence of cobalt addition. Changes in the electronic structure of

the solid when the cobalt solid solution is formed apparently change

the bonding enough to change the polarizability of these lattice

modes. Raman intensities depend directly upon the polarizability

tenser associated with each normal vibration.

However, the spectral selection rules do not change with cobalt

addition for these cathodically deposited materials. The spectra are

therefore consistent with a non-close packed structure, with the

cobalt going into solid solution on the nickel sites. We propose the

layer stacking is ABBCCA, similar to active mass.

Formation and cycling of the above materials produces an active

mass material that displays Raman spectra similar to those of cobaltfree, nickel electrode active mass (ref. 3). The fact that the active

mass spectral selection rules do not change upon cobalt addition

indicates that the cobalt is incorporated in solid solution, and that

a non-close packed structure is retained. The previous discussion of

187
active mass structure applies to cobalt containing active mass as

well.

Cobalt addition introduces a structure that is intermediate

between the _/_ and _/7 sto_chiometries. Both the Raman spectral

positions (550 and 474 cm -_, charged) and empirical formulae suggest

this (ref. 3). The recent cobalt-containing empirical formula

reported by Braconnier et al. can be recast in nonstoichiometric form

in equation 8 (ref. 14). The nickel deficit obtained is 0.19, falling

between the cobalt-free _/_ and _/7 values, as expected based upon

Raman spectral positions. It is apparent that the cobalt stabilizes

an increased cation deficit.

Ni0.57 Co0.24 K0.17 Na0.02 OO H0.92 Na0.08 (8)

STRUCTURAL AND CHEMICAL CHANGES

The Formation Process

Both cathodic-_ and active mass are nonst0ichiometric and nonclose packed. Therefore no change in layer stacking is required

during the formation process. If the cathodic-_ precursor contains a

significant nickel deficit, then the change in nickei nonstoichiometry

during formation can be minimal. For example, a cathodic-e (equation

6) would form an active mass as defined in equation 9 if no change in

nickel deficit occurred. A net loss in "water" is evident as well as

a change in oxidation state.

Ni0.75 (2H)0.25 OOH2.0"0.33 H20 + 0.25 OH- - ..... > (9)

Ni0.75 (2H) 0.25 OOH2. 0 + 0.58H20 + 0.25eThe formation of active_ass from a more stoichiometric precursor

would also involve creation of nickel vacancies. The final nickel

deficit characterizing an active mass need not be equal to nor

controlled by the deficit in the precursor-_. The formation process

can increase the nickel deficit to form more optimal active mass.

Transformation of a cathodic-e to active mass involves a change

in proton incorporation or bonding, as demonstrated by the loss of the

O-H stretching mode upon formation. This change in the spectrum is

consistent with our proposal that the cathodic-_ contains molecular

water which is removed or lost during formation.

A stoichiometric precursor with close packed stacking, such as _-

Ni(OH)2, must undergo major structural change upon formation into

nonsto[chiometric, non-close packed active mass. Both transformation

from close packed to non-close packed layer stacking and vacancy

creation are required. It is likely that layer shearing and defect

formation processes would slow the formation process for such starting

materials. Furthermore, shape and volume changes are expected to be

much greater for such a precursor. It is likely that these would

m

188
reduce the electrode mechanical integrity as well. We propose that

the reason that cathodically deposited materials are successful

precursors to active mass is intimately tied to the minimal structural

change required during the formation process.

Charge and Discharge Cycling and Transformation from _/_ to _/7 Cycles

Shifts in band position during charging and discharging of

cobalt-free active mass directly reflect changes in the structure and

bonding which occur as the nickel oxidation state varies (Table I).

The wavenumbers increase during charging and drop on discharge as

discussed earlier. The general charge-discharge reaction can be

written as shown in equation i0. K + may replace x of the protons to

give the _/7 cycle. If no overcharging occurs the structural changes

are simply changes in proton content and oxidation state. No change

in defect structure or layer stacking is required.

Nil_ x (nH)x OOH 2 + [l+x(n-l)] OH- - ..... > (i0)

Nil_ x H x OOHI. 0 + [l+x(n-l)] H20 + [l+x(n-l)] eOvercharging a 3_-material induces the 37-phase (ref. 2). As

seen in the nonstoichiometric formulae (Table II) the materials in the

two cycles display different nickel deficits. We propose that

overcharging induces the higher level of nonstoichiometry to

accommodate the higher oxidation state formed during overcharge. The

resultant point defect structure allows a greater change in oxidation

state, which provides the higher capacity observed for the _/7 cycle.

A significant amount of potassium is introduced when 37 is produced.

It is apparent that the potassium and the increased nickel deficit are

important to the stability of this structure.

Once formed, retention of the _/7 cycle might be expected.

However, further cycling will allow transformation back to the _/_

cycle with a concomitant reduction in nickel deficit and potassium

content. Furthermore, these are not the only structural changes

expected with this model. One mechanism for return to the _/_ cycle

involves aging of discharged-_ to disordered _-Ni(OH) 2 (as described

below), with subsequent cycling and re-formation bac_into the _/_

cycle. The intermediate structure is close packed (disordered 8-

Ni(OH)2 ) . Because this mechanism involves a change in layer-layer

stacking, density and volume changes are expected. Therefore, both

mechanical stability and capacity would be reduced. This is

consistent with experimental observations.

Aging in KOH

The product of KOH aging of a cathodic-_ was found to be

"deactivated-_" by Barnard e: al. (ref. 15). The Raman spectrum of

such an aged material shows that it is a nonstoichiometric,

disordered-_ compound (ref. 4). Barnard et al. observed a E-like x-ray

189
pattern as well, but did not differentiate it as a nonstoichiometric

material (ref. 9).

The Role of Cobalt Addition

The mechanisms by which cobalt has a beneficial effect on

electrode properties can now be discussed in terms of the structures

and structural changes already defined, cobalt is incorporated in

solid solution with nickel in cathodic-_ and active mass. All three

materials are non-close packed. Cobalt stabilizes nickel deficit

nonstoichiometry that is higher than for cobalt-free systems. This

structure enhances electrochemical properties, including a higher

nickel oxidation state.

Another role of the cobalt additive appears to be in stabilizing

a cathodically deposited precursor structure with enhanced

nonstoichiometry which is structurally similar to optimal active mass.

The very weak O-H stretching mode observed with cobalt addi£ion is a

measure of a water, proton and defect content more like that of active

mass. This will improve mechanical stability by minimizing the

density and volume changes during formation. The optimal

nonstoichiometry and point defect structure will also optimize the

electronic and protonic diffusion processes.

CONCLUSIONS

Optimal nickel active mass is non-close packed and

nonstoichiometric, as is cathodic-_. The f0rma£_on process transforms

precursor phases into this structure. Therefore, the precursor

nonstoichiometry and disorder, Or lack £hereof, influences the final

active mass structure and the rate of formation. Agfng to less active

material or incorporation of less than optimum nonstoichiometry

reduces capacity. The role of cobalt addition can be appreciated in

terms of structures favored or stabilized by the dopant.

The unique spectral signatures found for the various phases allow

differentiation of these materials and reflect significantstructural

variations. Most importantly, the layer stacking is clearly indicated

by Raman spectroscopic analysis. The nonstoichi0metry allows one to

understand the oxidation states and electrochemical properties. An

optimum active mass will exhibit an enhanced nickel deficit, providing

a maximum change in oxidation state, and will resist aging. How

formation of such a material can increase mechanical stability is

discussed in terms of a structurally Similar precursor-_.

This new structural understanding provides a basis upon which one

can study the subtle structural effects induced by empirical

parameters such as cobalt concentration or electrolyte concentration.

With such information the electrode structure can be tuned to provide

optimal electrochemical properties. It should be possible to relate

190
optimal structure of a precursor-_ and of formed active mass to

properties during cycling• Furthermore the potential of Raman

spectroscopy for quality control purposes appears bright because the

spectra vary qualitatively to mirror both major and subtle variations.

le

•

•

•

•

•

•

•

•

REFERENCES

Bode, H.; Dehmelt, K.; and Witte, J.: Zur Kenntnis der

Nickelhydroxidelektrode--I. Uber Das Nickel (II)-Hydroxidhydrat,

Electrochim. Acta, Vol. ii, 1966, pp. 1079-1087.

Bode, H.; Dehmelt, K.; and Witte, J.: Zur Kenntnis der

Nickelhydroxidelektrode. II. Uber die Oxydationsprodukte von

Nickel(II)-hydroxiden, Z. Anorg. Allg. Chem., Vol. 366, 1969, pp.

1-21.

Loyselle, P. L.; Karjala, P. J.; and Cornilsen, B. C.: A Point

Defect Model For Nickel Electrode Structures, Proc. Symp. on

Electrochemical and Thermal Modeling of Battery, Fuel Cell, and

Photoenergy Conversion Systems, J. R. Selman and H. C. Maru eds.,

Electrochemical Society, Pennington, NJ, 1986, pp. 114-121•

Cornilsen, B. C.; Karjala, P. J.; and Loyselle, P. L.: Structural

Models for Nickel Electrode Active Mass. J. Power Sources, Vol.

22, 1988, pp. 351-357.

Oliva, P.; Leonardi, J.; Laurent, J. F.; Delmas, C.; Braconnier,

J.; Figlarz, M.; Fievet, F.; and deGuibert: Review of the

Structure and the Electrochemistry of Nickel Hydroxides and

Oxyhydroxides., J. Power Sources, Vol. 8, 1982, pp. 229-255.

Shan, X.; Loyselle, P. L.; and Cornilsen, B. C.: Structural

Influence of Cobalt Addition to the Nickel Electrode, Abstract

79, 174th Electrochemical Society Meeting, Chicago, IL, October

14, 1988.

De Wolff, P. M.: Joint Committee on Powder Diffraction Standards

Card No. 14-117.

Szytula, A.; Murasik, A.; and Balanda, M.: Neutron Diffraction

Study of Ni(OH)2. Phys. Stat. Sol. B., Vol. 43, 1971, pp. 125-

128.

Barnard, R.; Randell, C. F.; and Tye, F. L.: Studies Concerning

the Ageing of Alpha and Beta Ni(OH)2 in Relation to NickelCadmium Cells. Power Sources 8, J. Thompson, ed., Academic

Press, London, 1981, pp.401-423.

191
i0.

Ii.

Loyselle, P. L.; Cornilsen, B. C.; Condrate, R. A.; and Phillips,

J. C.: X-ray Absorption Study of Beta Nickel Hydroxides, Abstract

77, 174th Electrochemical Society Meeting, Chicago, IL, October

14, 1988.

Nakamoto, K.: Infrared and Raman Spectra of Inorganic and

Coordination Compounds, John Wiley & Sons, New York, 1978.

12. Bode, H.: Joint Committee on Powder Diffraction Standards Card

No. 22-444.

13. McEwen, R. S.: Crystallographic Studies on Nickel Hydroxide and

the Higher Nickel Oxides. J. Phys. Chem., Vol. 75, 1971, pp.

1782-1789.

14.

15.

Delmas, C.; Braconnier, J. J.; Borthomieu, Y.; and Hagenmuller,

P.: New Families of Cobalt Substituted Nickel Oxyhydroxides and

Hydroxides Obtained by Soft Chemistry, Mat. Res. Bull., Vol. 22,

1987, pp. 741-751.

Barnard, R.; Randell, C. F.; and Tye, F. L.: Studies concerning

charged nickel hydroxide electrodes, i. Measurement of

reversible potentials, J. Appl. Electrochem., Vol. i0, 1980, pp.

109-125.

]92
-H

N

h

GI

.lab

Uo

m

,,¢m

m

o-_

-_1

.tJ .IJ •

_>

0_-P

1_ O

m ul .1..I

m,cI

m ,-I

-,-I

I 6-1

O

.r_

H

.Q

E-I

t_

-1-1

01

rA

m

.H

u

.,4 H

c.J

HHH

H

®=]

'O -,-I L

h_

o.

Itl -,-I

Ia ul_

U _-H

O

I_0

cOO

_)_)

0

tO

_0

f_

0

tO

_0

O

CO

In

CO

o_I

0 cO

t_

t_ ,_

(_1 co

c_ tO

tO') ,@

O_ tD

t_ r_

u'_ ,_

_D 0

I'-,

IFI ,_

'El

0

t/)

0

I_

,-I _ ,-i

e'l oO e"l

IFI I'- ,-I

193
Itl

0

I.i icl

-,-I ,-4

0

0

rJ

1_ °,-i

0t.i

M,I.)

.,4G)

_10

1_-,-i

O

o_

MO

HZ

[-t

-,-t O

-,-t Z

0 I

w

0

,-4

_ i o

_r_ I

O

-H r-t

o

_.1 O o

o

/_.tJ

O _

o

o

t_

A

O

,4

iln

o

o

.,-i i

o

IN

"o _

. 0 0

I_I 0 0

0 _

0 0 0

l.rl

It) -H -,-I

c;

-,-I

("q

o 0

_ o o

_ o

o _;

o

d =o

I>

_ o

._.t ° _

d

-,-i

113 113 _ Iln

_ O

0 I"- OO I_ I_

0 0 0 0 0

0 0 0 0 0

I

0

--4

E_

4_

194
44

0

l/J

-,-I 0

-IU

m'Ci

-,-I G)

_m

-IJl_

U

tllUl

elO

01.4

_/1 ,-4

.i-i O _1

(/1

-,-t O

r-4

Or0

°_

H_

t-!-,'4

H_

O

,1:1

r-4 O

O

O

• i o

;> 14 11,

.4 _ ro

_, o .._ o

IFI

• 0 -,4 0

• 0 -_ 0

u_

• 0 -M 0

e_ C _ C

0

]95
196
N90-20473

IMPEDANCE STUDIES OF Ni/Cd AND Ni/H 2 CELLS USING THE CELL CASE

AS A REFERENCE ELECTRODE

Margaret A. Reid

NASA Lewis Research Center

Cleveland, Ohio 44135

Impedance measurements have been made on several Ni/Cd and Ni/H 2

flightweight cells using the case as a reference electrode. For these

measurements the voltage of the case with respect to the anode or cathode

is unimportant provided that it remains stable during the measurement of

the impedance. In the cells measured so far, the voltages of the cell

cases with respect to the individual electrodes differ from cell to cell

even at the same overall cell voltage, but they remains stable with time.

The measurements can thus be used to separate the cell impedance into the

contributions of each electrode, allowing improved diagnosis of cell

problems.

INTRODUCTION

Many impedance studies have been carried out on Ni electrodes and

Ni/Cd and Ni/H2 cells. Some studies have been made while the cells or

electrodes were being discharged (refs. I-2), others at various open

circuit voltages (refs. 3-6). The impedances have been found to be

strongly dependent on depths of discharge and/or voltage and on the

procedure used to obtain the measurements. The studies have been similar

qualitatively but not quantitatively.

In order for impedance to be used as a diagnostic too] for individual

cells, more reproducible measurements must be obtained. It would also be

desirable to be able to separate the contributions of the individual

electrodes, especially for Ni/Cd cells. With Ni/H 2 cells one can almost

always assume that any problems are due to the Ni electrode or separator

and that the hydrogen electrode behaves reasonably reversibly. In this

case the cell impedance is essentially that of the nickel electrode plus

the ohmic contributions from the separator and electrolyte. However, with

Ni/Cd cells, the Cd electrode can sometimes be the cause of poor cell

performance although it usually causes fewer problems than the Ni

electrode. For a sealed cell with no reference electrode, there is

currently no satisfactory way of determining which electrode is behaving

poorly without opening the cell and performing postmortem analyses.

This paper discusses preliminary measurements of impedance on

flightweight Ni/Cd and Ni/H 2 cells using the case as a reference electrode

in order to determine the impedances of each individual electrode.

Although measurements so far have been limited, the method appears valid.

We are reporting it at this time because of the urgent need for better

diagnostic procedures for space-qualified Ni/Cd and Ni/H 2 cells.

]97
EXPERIMENTALPROCEDURE

Impedancemeasurementswere madeusing the Solartron 1250 Frequency

ResponseAnalyzer coupled with a Solartron 1286Electrochemical Interface.

Measurementswere madeon two flightweight 50 AHNi/Cd cells from General

Electric Corporation (now Gates) which had undergone acceptance testing

several years ago and which had subsequently been stored in the shortcircuit condition. Onewas from a lot which had Cd electrodes which had

been determined to be faulty from testing of other cells of the lot, the

other was believed to have a good Cd electrode. At the time of testing,

the cells had been stored for about six years. They were measuredin the

as-received condition, partially charged at I A for 5 hours, discharged

at I A to 0.I V, where they were measuredagain. They were then charged

at 5 A (C/lO) for ]6 hours, discharged at 10 A (C/5) to 0.i V and

equilibrated at 1.275V. The impedances were then measured at room

temperature from lO0 Hz to 0.0025 Hz with a RMSsignal of I mV(peak to

peak signal of 2.828 mV). .........

The Ni/H2 cell studied was a 50 AHflightweight cell from Hugheswhich

had completed acceptance tests and been cycled 900 reaT-time LEOcycles,

followed by three years on open-circuit. The cell was charged to 1.275

V and equilibrated overnight. Impedance measurements were then taken of

the total cell and the individual electrodes. The cell was then charged

several times at C/I0 for 18 hours and discharged at C/2 to 1.00 V.

Following this it was cycled at 604 DOD for three LEO cycles (55 minutes

charge, 35 minutes discharge, 5_overcharge). Impedance measurements were

taken again after equilibration of the cell at 1.275 V. The same

parameters for the impedance measurements were used as with the Ni/Cd

cells.

RESULTS AND DISCUSSION

A simplified circuit diagram lforan electrochemical celi is shown in

Figure I. A real cell usually has additional circuit elements including

inductance due to the physical arrangement of the plates, modifications

due to the porous nature of the electrodes, and adsorption capacitances

and resistances. However, this simplified circuit diagram serves as a

good first approximation. If the impedance of one of the electrodes is

several orders of magnitude less than that of the other as is usually the

case with Ni/H 2 cells, one of the RC loops can be neglected, and the

impedance uf the cell can be treated as essentially that of the electrode

with the higher impedance. The impedance data can be plotted in a number

of different ways (ref.7). Bode plots (the log of the magnitude of the

impedance vs the log of the frequency) are useful for comparing data that

vary by orders of magnitude. Nyquist or complex plane plots (the real or

in-phase component of the impedance vs the imaginary or out-of-phase

component) are helpful in determining the ohmic and kinetic resistances

and double-layer capacitances, and the slope of the Warburg plot (either

the real or imaginary component vs the inverse square root of the angular

frequency) gives information about the diffusion process (a large slope

corresponds to slow diffusion). Figure 2 shows a typical Nyquist plot of

a well-behaved Ni/H2 cell at a voltage where the kinetic resistance is

low. The impedance in this case is essentially that of the Ni electrode.

The left intercept with the x-axis is the ohmic resistance, and the

198
diameter of the semicircle gives the kinetic resistance. The right-hand

portion of the curve is caused by slow diffusion processes. Similar plots

are obtained with Ni/Cd cells (refs. 4-6), and the bulk of the impedance

is also usually due to the Ni electrode. At low voltages, the diffusion

portion of the curve may merge with the kinetic portion.

If the impedance of the cell varies from that of the typical cell,

or if there is reason to suspect problems with the Cd or H2 electrodes, it

would be desirable to separate the impedance into the contributions from

each individual electrode. This study was undertaken in order to see if

the use of the cell case would enable such a separation.

In studying the impedances of cells containing Ni electrodes, the

problem arises as to the best voltage or voltages for carrying out the

measurements. It has been found that the impedances of Ni electrodes in

the charged state (i.e., at higher voltages) are very low and do not vary

much from one electrode to another or from one manufacturer to another.

However, at low voltages corresponding to almost completely discharged

electrodes, impedances rise by several orders of magnitude. Figure 3 is

the Bode diagram for the magnitude of the impedance vs frequency for a

typical electrode showing the changes with voltage as well as with

frequency. Figures 4 and 5 show how the kinetic resistance and the

Warburg slope change abruptly by several orders of magnitude as the

voltage is lowered. At very low voltages (cell voltages of less than

1.000 V or electrode voltages below about 0.050 V vs the Hg/HgO electrode)

the impedances are very large, and there is a great deal of scatter. We

have found that equilibrium cell voltages of 1.150 to 1.275 V

(corresponding to voltages of the Ni electrode vs the Hg/HgO electrode of

about 0.200 to 0.325 V) give sufficient differentiation of the impedances

with less scatter than at lower voltages. The measurements reported below

were taken at cell voltages of 1.275 V.

We have also found significant differences at low voltages between

electrodes from different manufacturers in addition to small differences

between individual electrodes from the same manufacturer. This is not

important if one is comparing cells from a single manufacturer, but must

be taken into account if comparisons are made between electrodes or cells

from different sources.

Recent studies in our laboratories have also indicated that

measurements of impedances of the Ni electrode are more reproducible if

the cells or electrodes are equilibrated at the voltage of interest.

Equilibration times required vary from a few minutes for highly charged

electrodes to several hours or overnight or longer for electrodes at low

voltages. The voltage range described above allows reasonable

equilibration times of 15 min to several hours depending of the electrode

and the magnitude of the change from the initial voltage to the

equilibrium voltage desired.

The voltage of the Cd electrode is reasonably stable, and it can be

used as a reference electrode although it is less satisfactory than the

Hg/HgO electrode (ref. 8). If the electrode is partially charged so that

there is a reasonable amount of both metallic Cd and Cd(OH)2 present, the

potential should be quite close to the theoretical value (Eu=O.809 V vs

199
SHE). Thus if the voltage of the cell case vs the Cd electrode is fairly

constant, we can assume that the cell case voltage is not changing

significantly.

The voltages of the cell cases with respect to the Cd and Ni

electrodes at various times are given in Table I. We see that the

voltages of the case vs the Cd electrodes are quite stable, especially

over short periods of time, while the voltage of the Ni electrode vs the

case varies gradually with time. Although it is only necessary that the

case voltage remain constant during the impedance measurement, it is

interesting to speculate as to what electrode reaction is determining the

potential of the case. The standard voltages of several possible couples

are presented in Figure 6 along with the couples for the Ni and Cd

electrode reactions, the Hg/HgO reference electrode and the case voltages

for the two Ni/Cd cells studied

Flightweight Ni/Cd cells have stainless steel cases. If we compare

the data in Table I with the values in Figure 6, we see that the potential

of the case with respect to Cd for cell L]3-012 is somewhat higher than

that of the Fe(OH)3/Fe(OH) 2 potential, that of the other cell is not close

to that of any obvious couple. There are no other metal/metal oxide

potentials in this voltage range that could reasonably be considered as

the source of this potential. Thus the potential-determining reaction

cannot be determined at this tlme, but sincethe potential is stable for

a reasonable time period, determination of the impedance using the case

as the reference is feasible The potential may be a mixed potential Or

it may be affected by organic components from degradation of the

separator.

Impedance measurements were made on the complete cell and on each of

the electrodes individually using the case as a reference. Figure 7 shows

the arrangement of the leads for each measurement. The voltage between

the two reference electrodes was maintained constant during the

measurement. The experiments were interrupted several times during the

course of the measurements of the individual electrode impedances to see

if the ceTi Voltage and the voitage of the electrode which was not being

measured had changed. If the cell voltage had fallen (in no case this was

more than 3-4 mV}, the cell was again brought to the initial voltage

before the measurements were continued.

The results are plotted either as Bode plots or Nyquist plots. The

Nyquist measurements were corrected for ohmic resistance since it was not

known initially whether the sum of the ohmic resistances between the case

and the individual electrodes would equal the ohmic resistance between the

two electrodes. (The results of the measurements subsequently showed that

this was unnecessary, that each of the resistances between the case and

the individual electrodes was half that of the ohmic resistance of the

complete cell). The corrected impedance for the two electrodes was added

together at each frequency to give a calculated impedance for the complete

cell. This calculated impedance was then compared to the experimental

impedance for the cell (which had also been corrected for ohmic

resistance). The values were essentially the same except at low

frequencies where the impedances are higher and the scatter greater,

indicating that the method is valid.

2OO

=

m
Figure 8a is the Nyquist plot for the impedances in cell L13-0]2, the

cell which was presumed to have a poor Cd electrode. It can be seen that

the sum of the impedances for the individual electrodes at each frequency

is equal to the measured impedances for the total cell. The same behavior

is shown in Figure 8b for cell L07-220, from a lot with "good" Cd

electrodes.

Figure 9a compares the impedances of the "good" and "bad" Cd

electrodes using the Bode representation. The impedance of the "bad" Cd

electrode in cell L13-012 is higher at low frequencies (diffusion

controlled region) by half an order of magnitude. Figure 9b is the

comparable plot for the Ni electrodes. In this case the low frequency

impedance of the Ni electrode in cell L13-012 is lower than that in cell

L12-220. The Nyquist plots are shown in Figure 10.

The values of the kinetic resistances, Warburg slopes, and electrode

capacitances that can be obtained by analysis of the data using the

circuit diagram for an individual electrode are not reported, since we

believe that the cells had not been cycled sufficiently after the

open-circuit or short-circuited storage to obtain reproducible results.

The data are presented primarily to describe the method and indicate its

probable validity. However, we feel that the results are qualitatively

correct.

The data for the Ni/H 2 cell after cycling are presented in Figure 11.

The Bode representation is used since the nickel electrode impedance and

the hydrogen electrode impedance differ by about two orders of magnitude.

The impedance parameters are given in Table II.

In both sets of measurements of the Ni/H 2 cell the voltage of the cell

case was not as stable as with the Ni/Cd cells. Sometimes the voltage was

close to the Ni potential, at other times close to potential of the H2

electrode. During the measurements of the impedance of the Ni electrode

the case potential remained stable, i.e., the cell voltage after the

measurement was essentially the same as before, and no appreciable DC

current was noticed. However, during the measurement of the impedance of

the hydrogen electrode the cell voltage changed by several hundred mV and

a large DC current was noticed. We suspect that the instability of the

case voltage may be due to the difference in the case material from that

of the Ni/Cd cells. The case of the Ni/H 2 flightweight cell is made of

Inconel, while that of the Ni/Cd cells is steel. The voltages of the case

vs Ni in nine boilerplate Ni/H 2 cells housed in standard steel pressure

vessels are in the same range as those of the two Ni/Cd cells described

above and seem to be as stable as in the Ni/Cd cells. The difference is

thus probably due to the case material rather than the difference in the

negative electrode or separator. Despite the changes in voltage of the

case in the Ni/H 2 cell, the results are believed to be qualitatively

correct.

201
CONCLUSIONS

Preliminary experiments indicate that the cell case can be used as

a reference for Ni/Cd cells. The case can also probably be used as a

reference for Ni/H2 cells, although the potential of the case does not

appear to be as stable as with the Ni/Cd cells. Thesemeasurementscan

be used to help determine which electrode is responsible for loss of

performance. Weare now cycling these Ni/Cd cells to determine changes

in impedancesof the individual electrodes with cycle life. Wehope to

makemoremeasurementson a variety of Ni/Cd and Ni/H2 cells from different

lots to see if the cell and electrode impedancescan be correlated with

cycle life.

202
°

,

,

REFERENCES

Zimmerman, A., Martinelli, M.R°, Janecki, M.C., and

Badcock, C.C., J. Electrochem. Soc., 129 (1982), 289.

Barton, R.T., Hughes, M., Karunathilaka, S.A.G.R., and

Hampson, N.A., J. Appl. Electrochem., 15 (1985) 399.

Lenhart, S.J., Macdonald, D.D., and Pound, B.G., J.

Electrochem. Soc., 135 (1988) 1063.

4. Haak, R., Ogden, C., and Tench, D., J. Power Sources, 12

(1984) 289.

5. Comtat, M., Delorme, P., Baron, F., and Capel, A., Proc.

ESA Session at 16th Annual IEEE PESC (ESA SP-230, May

1985.)

6. Reid, M.A., Extended Abstracts of the 174th

Electrochemical Society Meeting, Chicago, IL, Oct. 1988,

122.

r

7. Bard, A.J., and Faulkner, L.R., Electrochemical Methods.

Wiley, (1980) 316.

8. Salkind, A.J., and Bruins, P.F., J. Electrochem. Soc.,

109, (1962) 356.

203
1

Z

W

1

l w

Fig. i. Simplified circuit diagram for an electrochemical cell.

Fig. 2.

0.014

0.012

In

E O.OLO

o

0.008

o

c-

,m

im

o 0.006 E

N

0.004

0.002

0

0

V = 1.200

I I I I I

0.002 0.004 0.006 0.008 o.oio

Z reol, ohms

Nyquist plot of impedance of typical 65 AH Ni/H 2 cell

being tested for Space Station Freedom.

204
U_

nO

LJ

O

Z

,<

E3

W

n

O

o

..J

1

0

-1

-2

-3

VOLTAGE vs Hg/HgO

• 0.410 V

0.330 V

o 0.250 V

A 0.170 V

X 0.125 V

v -0.927 V

I I I I I

-1 1 3

LOG FREQUENCY, Hz

Fig. 3. Dependence of impedance on voltage for a Gates electrode,

42.9 cm _, 31% KOH, after 12 cycles.

e4

a

o

W

-I

-2

VOLTAGE VS Hg/HgO , MV

Fig. 4. Changes of kinetic resistance with voltage,

electrode, normalized.

Gates

205
2

%

b/

U

O

J

O

q

-2

-3 I ! I ! ! l_ ! !

VOLTAGE VS Hg/HgO*

Fig. 5. Changes of Warburg slope with voltage, Gates electrode,

normalized.

-i.0 -0.5 0:0 0.s

, I , , , _, I , , , , I

OH-/ rFe(OH)3/Fe(OH) ' 2 02/H02 HgO H NiOOH/Ni(O 2

Cd(OH)2/Cd Co(OH)3/Co(OH) 2

Ni,Co/Ni( OH )2'Co( OH )2

Fig. 6. Standard reduction potentials in basic solution and

potentials of cell cases for Ni/Cd cells L07 and LI3.

m

B

=

_=

=

@

IMPEDANCE OF COMPLETE CELL

Fig. 7. Arrangement of leads for impedance measurements.

206

•
O

O

_E

N

2"

2.8-

2.6

2.4

2.2-

2-

1.8-

1.6-

1.4.

1.2.

1-

0.8-

0.6-

0.4-

0.2-

0

0

I CELL,EXP

Ni vs CASE

0 Cd vs CASE

A CELL.,CALC

O

C

,--

O

._E

N

0.2, 0.4, 0.6, 0.8, ,1 12: 14: 0.2 0.4 0.6 0.8

Z real Z rear

2.8- z CELL,EXP

2.6 _ Ni vs CASE

, Cd v-CASE ,_.,)'

2.4 _ CELL,CALC /_.

2,2- "

1.8-

16- .,

114

0.6

0.4

O,

u_p" I I I I I I "1' I

0 1 1.2 1.4

Fig. 8. Comparison of sums of individual electrode impedances of

Ni/Cd cells with experimental cell impedances, Nyquist

plot. Cell voltage = 1.275 V. a) Cell LI3, b) Cell

L07. (Scale of axes is unequal in order to show data

more clearly.)

• LOT

O * L13

!-,

,,

-_ -i__-_--_ -_ F---

LOGFREQUENCY.Hz

• 1.07

u_ 0

i:

- g-

-_ 22 2_ ' 6 ' i

LOGFREQUENCY.Hz

Fig. 9. Impedances of individual electrodes in cells L07 and LI3,

Bode plot. Cell voltage = 1.275 V. a) Cd electrodes,

b) Ni electrodes.

207
2-

I.B1.6-

1.4

°

1.2-

,_c

.E

N 0.8

0,6-

0.4

0.2-

0_

0

Fig. i0.

II LO7.good electrode

L13, bad electrode

02 04 06

Z reol

'o18

_E

N

2-

1,8i

1.6"

1.4_

1.2

!

O.8.J

0.6

0.4

0°2

0

e L07 /

0.2 0.4 0.6 0.8

Z reol

impedances of individual electrodes of Celis L07 and LI3,

Nyquist plot. Cell voltage = 1.275 V. a) Cd electrodes,

b) Ni electrodes. (Scale of axes is unequal in order to

show data more clearly.)

11

I

O

z

o -2

<

N"

o

o

d

I

• H vs CASE

* Ni vs CASE

NI vs H, CALC.

" Ni vs H. EXP.

X\

-- ___ - _

[ 'I _ '

!

- I -

LOG FREQUENCY, HZ

Fig. ii. Comparison of sums of individual electrode impedances in

50 AH Ni/H 2 cell with experimental cell impedance, Bode

plot. CelI voltage = 1.275 V. =

208
Table I. VOLTAGES OF

ELECTRODES.

CASES OF Ni/Cd CELLS VS INDIVIDUAL

Cell L13-012 At Time of

Impedance

Overnight

Stand

After

Recharge

Two Months

OCV

Cell Voltage 1,275 1,1 81 1.275 0.243

Case vs Cd 0.350 0.350 0.349 0.318

Ni vs Case 0.925 0,831 0.926 -0.076

Cell L07-275

Cell Voltage 1.275

Case vs Cd 0.648

Ni vs Case 0.627

0,413

0.6,31

-0.218

Table II. IMPEDANCE PARAMETERS FOR Ni/H 2 50 AH CELL.

COMPLETE CELL

NI vs CASE

H vs CASE

Rohm P_in SLOPE CAPACITANCE

(mOhm) (mOhm) (Forads)

4.00 - .0063 -

1.96 - .0041

2.05 - .000056 -

209

N90-'20474

THE APPLICATION OF ELECTROCHEMICAL IMPEDANCE SPECTROSCOPY FOR

CHARACTERIZING THE DEGRADATION OF Ni(OH)2/NiOOH ELECTRODES

D.D. Macdonald and B.G. Pound

SRI International

Menlo Park, California 94025

and

S.J. Lenhart

Ford Aerospace

Palo Alto, California 94303

Electrochemical impedance spectra of rolled and bonded and sintered

porous nickel battery electrodes were recorded periodically during

charge/discharge cycling in concentrated KOH solution at various

temperatures. A transmission line model (TLM) was adopted to represent the

impedance of the porous electrodes, and various model parameters were adjusted

in a curve fitting routine to reproduce the experimental impedances.

Degradation processes for rolled and bonded electrodes were deduced from

changes in model parameters with electrode cycling time. In developing the

TLM, impedance spectra of planar (non-porous) electrodes were used to

represent the pore wall and backing plate interfacial impedances. These data

were measured over a range of potentials and temperatures, and an equivalent

circuit model was adopted to represent the planar electrode data. Cyclic

voltammetry was used to study the characteristics of the oxygen evolution

reaction on planar nickel electrodes during charging, since oxygen evolution

can affect battery electrode charging efficiency and ultimately electrode

cycle life if the overpotential for oxygen evolution is sufficiently low.

Transmission line modeling results suggest that porous rolled and bonded

nickel electrodes undergo restructuring during charge/discharge cycling prior

to failure. The average pore length and the number of active pores decreases

during cycling, while the average solid phase resistivity increases. The

average solution phase resistivity remains relatively constant during cycling,

and the total porous electrode impedance is relatively insensitive to the

solution/backing plate interfacial impedance.

INTRODUCTION

Porous nickel electrodes are used in a number of secondary alkaline

battery systems, including nickel-iron, nickel-zinc, nickel-hydrogen and

nickel-cadmium cells. Each of these batteries must ultimately meet several

performance criteria: high specific power, high specific energy, low cost, and

long cycle life. At present, the viability of these batteries is often

21]

PRECEDING PAGE BLANK NOT FILMED
limited by the degradation of the electrode materials. In some cases, the

nickel plate is life-limiting (Ref. i).

A number of irreversible degradation processes affect the performance of

porous nickel battery electrodes. Like all porous electrodes, nickel plates

can exhibit electrolyte exhaustion within the pores leading to mass transport

and ohmic overpotential losses that reduce cell power. Faradaic efficiency

losses from cycle-and temperature-dependent parasitic processes, such as

oxygen evolution, can reduce charging efficiency. Also, the structural

integrity of porous nickel electrodes frequently is inadequate to endure the

mechanical stresses that arise during charge/discharge cycling. Resistive

degradation of substrate particle-particle bonds can result from these

stresses, and/or the active material may progressively separate from the

current collector further reducing the performance of the electrode on cyclic

charging and discharging.

In this paper, we report a study of the degradation of porous nickel

battery electrodes in alkaline media upon cyclic charging/discharging.

Electrochemical impedance spectroscopy is used as the principal experimental

tool. Impedance studies of both pianar and porous nickel battery electrodes

in alkaline solutions have been published previously, but much of this work

was restricted to relatively narrow frequency ranges because of limitations

with experimental instrumentation (Ref. 2-5). Also, some investigators report

impedance data for the total cell rather than the individual electrodes (Ref.

6,7), while other studies have dealt with electrodeposited (thick) oxide films

(Ref. 8).

TRANSMISSION LiNE MODEL

An understanding of how the properties of porous nickel electrodes are

altered during cycling is developed in this study by adopting a transmission

line model (Ref. 9,10) for the impedance of the porous mass. The model is

adopted from Lenhart, Chao, and Macdonald (Ref. Ii) and Park and Macdonald

(Ref. 12), and differs from classical TLMs in two ways. First, the model used

here recognizes the finite thickness of a practical battery electrode.

Accordingly, the electrochemical behavior of the porous mass will be partly

determined by processes that occur at the base of the pore between the current

collector (backing plate) and the solution (impedance Z', Figure I), provided

that the frequency is sufficiently low that the penetration depth of the AC

wave is of the same order as the thickness of the porous mass. Secondly, the

model assumes a finite resistance for the active solid phase in order to

account for the resistive degradation of particle-particle contacts caused by

internal stresses.

As with most porous electrode models developed to date, several

simplifying assumptions are made in order to render the mathematics tractable.

Thus, the pores are assumed to be parallel right cylinders, and any radial and

axial eiectrolyte concentration gradients within the pores are neglected.

Furthermore, average pore electrolyte and solid phase resistances are used. A

uniform layer of active material is assumed to line the walls of the pores,

2]2
and charge storage processes along the walls are represented by a positionindependent interfacial impedanceZ. In this work, an equivalent circuit

representing the interfacial impedance, Z, and the backing plate impedance,

Z', are deduced from planar nickel electrode impedances.

The mathematical details of the modified transmission line model have

been described in previous publications (Ref. 11,12) and are discussed only.

briefly here. The equivalent electrical circuit for a single pore in

discretized form is shownin Figure 2, in which Rm and Rs represent the

resistance of the solid current-carrying phase per unit pore length (ohm/cm),

and the resistance of the solution phase per unit pore length (ohm/cm),

respectively. The i_terfacial impedance, Z, is a specific impedanceper unit

pore length, (ohmcm_/cm), so Z/dx has units of ohms. The current collector

or backing plate impedanceZ' has units of ohms, and is assumedto be

independent of pore length.

Current and potential distributions within the porous system, and the

total impedance, were derived (Ref. 11,12) by application of circuit analysis

equations to a typical discrete unit. The total impedanceof n onedimensional parallel pores was found to be:

_-_RmRs2 271/2RR + 7 1/2(R2 + R2) C + 6 R2S_ - I m s m s s

ZT nlR +R + (i)

m s 1/2 1/2

7 (Rm + Rs ) (7 S + 6 C)

where

R R R + R

m + s m s C cosh (7 1/22), and S sinh (7 1/22). 7 - , 6 - - -

If A isZthe total pr z'jected' area of the porous electrode, and (i-8) is the

fraction of that area occupied by pores, then 8A/n is the film area per pore,

and (l-8)A/n is the average pore area per pore. The solution phase resistance

per pore becomes Psn2/(l-8)A, and the resistance of the current-carrying solid

phase is Pmn2/eA, where Ps and Pm are the resistivities (ohm cm) of the

solution and solid phases, respectively, and 2 is the pore length. The

resistances R s and Rm in Figure 2 are therefore

Rs - Ps n/(l-e)A (ohm/cm) (2)

Rm - Pm n/SA (ohm/cm) (3)

If the specific impedances (ohm cm 2) of the pore wall and pore base are Z

and Z', respectively, then the impedance of the pore wall and pore base per

pore are Zw/2_rl and Z_/_r 2, where r is the average pore radius. Since the

average pore area is (_-8)A/n = =r 2, the average pore radius is given as

r = [(1-0) A]I/2

_n

(4)

2]3
and the impedance per unit pore length as:

Z 2 Z

Z = 2_r_ " 2 _ A (l-e)

Similarly, the backing plate or current collector impedance Z' per pore is

found to be

Zb Zbn

Z# J -- z

2 (I-8)A (6)

_r

The above expressions for Rm, Rs, Z and Z' are used in Equation i, which now

describes the impedance of a three dimensional porous electrode. The

expressions for Z and Z' are determined from planar electrode impedances, as

discussed in the Results Section.

==

EXPERIMENTAL

Test Cell

A three electrode cylindrical PTFE cell was used for all experiments with

the working electrode positioned horizontally near the bottom of the cell. A

platinum counter electrode and a Hg/HgO reference electrode were positioned

over the working electrodes. The cell provided input ports for the _

electrolyte solution, for high purity argon gas purging, and for a PTFE coated

copper/constantan thermocouple. High purity argon purging gas was

deoxygenated in two zinc/vanadyl sulfate gas washing bottles. An 8 molal KOH

electrolyte solution with 1% LiOH was used for all experiments and was

prepared from reagent grade KOH and LiOH in double distilled, deionized

water. The small LiOH addition was made to conform with other previously

reportedbattery cycling experiments. Lithium additions are usually regarded

as beneficial to porous electrode performance (Ref. 13,14,15,) although recent

experiments (Ref. 16) showed little effect on capacity during cycling.

All potentials reported here are relative to the Hg/HgO reference

electrode. A paste of Hg/HgO was inserted in a PTFE container above the

solution, and contact with the electrolyte was provided by cotton fibers in a

PTFE capillary. No liquid junction correction was required with this

arrangement, since the reference and working electrodes were in contact with

the same electrolyte. From the reaction

HgO + H20 + 2e Z Hg + 20H (7)

the potential of the Hg/HgO reference in 8 molal KOH was calculated to be

=___

i

214

z
-0.008V (SHE) using the following values for the activity of water and OH-:

aH20(298°K ,w 8m KOH) - 0.5545, E ° - 0.0984 V and aOH- - (mOH-)

V, - mKOH? * = (Smolal)(5.902) from Pound et al. (Ref. 17).

Three kinds of nickel working electrodes were used: a planar nickel

electrode, rolled and bonded porous electrodes, and sintered porous

electrodes. The planar nickel specimen was cut from a rod of 99.5_ nickel. It

was polished to a 0.05 micron alumina powder finish and was rinsed with

distilled water. Typically less than I0 minutes elapsed between polishing and

polarizing the sample, and only a few seconds elapsed between solution contact

and polarization.

The porous electrodes used in this study were prepared by commercial

electrode fabricators. The active material in the rolled and bonded electrodes

was supported by a PTFE "web" making up 1 w/o of the total electrode material;

the remainder being 30_ graphite, I_ cobalt hydroxide, and hydrated nickel

hydroxide. The graphite served as the current carrier to the backing plate,

and the cobalt was added to increase capacity during cycling (Ref. 18). A

capacity of 0.29 A hr/gm was reported based on a one electron transfer from

nickel hydroxide to nickel oxyhydroxide. The structural features of the

sintered electrodes were very different from the rolled and bonded

electrodes. For the former, nickel powder was sintered to a nickel wire mesh,

and NiOOH was chemically deposited in the pores. The sintered nickel metal

(and not graphite) carried the current to/from _he wire mesh, which served as

a backing plate. The capacity was 0.015 A hr/cm projected (flat) area.

Experiments were performed at temperatures ranging from 0 to 100°C.

temperature was controlled to within , 2°C as indicated by a thermocouple

inside the test cell.

The

Cyclic Voltammetry

Only planar nickel electrodes were studied by cyclic voltammetry.

Freshly polished nickel electrodes were inserted into the cell and, on contact

with the electrolyte, were polarized to -850 mV. The solution and cell were

then heated or cooled to the desired temperature. After temperature

stabilization, a triangular potential/time perturbation was applied to the

cell via a coupled function generator and potentiostat.

Various sweep rates from 1 to I00 mV/s were employed, and the potential

was swept from below the hydrogen evolution potential to well above the oxygen

evolution potential before reversing the sweep direction. At a given

temperature, E/I traces were first recorded at i00 mV/s, then at progressively

lower sweep rates. After the lowest sweep rate voltammogram was recorded, E/I

traces were recorded at consecutively higher sweep rates up to I00 mV/s. The

volta_mograms reported here were reproducible to within , 3 mV and , 0.5

mA/cm_on the potential and current scales, respectively.

215
Electrochemical Impedance Spectroscopy

Impedance data were recorded with either a Solartron 1172 or 1250

Frequency Response Analyzer. For all impedance measurements, the Solartron

sine wave output was superimposed on an applied DC bias from a Princeton

Applied Research Model 173 potentiostat. Solartron potential and current

input leads were taken directly from the cell and not from the potentiostat

electrometer and current output jacks. A unity gain voltage follower based on

an AD 521J operational amplifier was placed between the cell and the potential

input of the Solartron to avoid polarizing the reference electrode. The

amplifier had a differential input impedance of 3 X I0° ohms, and a flat

frequency response (f I_) at unity gain to 75 kHz. The voltage follower was

accurate to 0.I mV DC relative to a digital voltmeter.

Planar electrode impedance spectra were recorded over a range of DC

potentials. Impedances were usually measured sequentially without repolishing

the electrode between measurements. Electrodes were first polarized for two

hours at the lowest potential of a given measurement sequence (typically -150

mV). The impedance spectrum was recorded, followed by a potential step

(usually i00 mV) to the next highest potential. After one hour at the higher

potential, another impedance spectrum was recorded. This procedure was

normally repeated up to about 500 or 600 mV.

Porous electrode impedances were recorded in th_ fully discharged

condition (0 mV). They were recorded periodically after selected numbers of

charge/discharge cycles. The cycling process is described below, and one hour

elapsed at constant potential (0 mV) before impedance spectra were recorded.

Charge/Dischar£e Cycling

Porous electrodes were cycled at constant current using an ECO Model 545

Galvanostat/Electrometer. Various charging currents were used, but the

electrodes were always discharged at twice the charge rate. They were usually

charged to I00_ of rated capacity, and were fully discharged (I00_ DOD) on

each cycle. Four or five "conditioning" cycles were completed before

impedance data were recorded.

The galvanostat provided for automatic current reversal at selected

potentiais by presetting front panel Poten_i6meters. However, in most cases,

it was necessary to use constant charging and discharging times. Two timers

were used to control switches connected to charge and discharge trigger inputs

on the ga!vanostat. When charging efficiency was less than i00_, it was

necessary tO stop the discharging current before the end of the set discharge

time. A voltage comparator based on an LM 311 amplifier was constructed and

included in the timer circuitry to stop the discharging current prematurely at

any selected potential until the next charging cycle started.

v;

i

z

216
RESULTS AND DISCUSSION

Planar Electrodes - Cyclic Voltammetry

Cyclic voltammetry was used to determine the extent of oxygen evolution

during nickel hydroxide oxidation. Oxygen evolution is a parasitic reaction

during charging of nickel battery electrodes, and oxygen gas bubble formation

may contribute to electrode degradation by generating internal stresses within

the electrode pores. The large KOH concentration and elevated temperatures

used in this study serve to enhance oxygen evolution by decreasing the

overpotential.

For most cycling experiments, only one anodic oxidation peak, appearing

at about 500mV, was recorded prior to oxygen evolution. Similarly, only one

oxyhydroxide reduction peak at about 300mV was observed on the reverse sweep.

Similar voltammograms have been reported for nickel in various alkaline

solutions (Ref. 14,19).

At the highest sweep rate of 100mV/s, no steady state voltammogram was

observed even after cycling continuously for over 19 hours (Figure 3). Both

the anodic peak currents and anodic charge were found to increase steadily

with time, but at a decreasing rate. Cathodic peak currents and the

associated cathodic charge were difficult to determine since the cathodic

current base line was obscured by the oxygen evolution current. However,

minimum and maximum values for the cathodic parameters were estimated. For

the first few tens of cycles, results indicate that more charge is consumed

during hydroxide formation on the cathodic sweep than is liberated during

oxyhydroxide formation on the preceding anodic sweep (qc>qa) . This is

possible if oxygen becomes trapped within the film or does not desorb from the

film/electrolyte surface rapidly, and is reduced during the subsequent

cathodic sweep (Ref. 20). After several additional cycles, the anodic charge

becomes larger than the cathodic charge (qa>qc) . This suggests that athe

nickel substrate oxidizes during cycling and possibly that some film

dissolution occurs.

Cycling time at i00 mV/s had only a minor effect on the proximity of the

hydroxide oxidation peak to the oxygen evolution line at ambient

temperature. The anodic peak potentials initially decreased and then

increased with cycling time, but overall the changes were small, as indicated

in Figure 3.

Anodic peak potentials were closer to oxygen evolution curves at higher

temperatures. Figure 4 shows voltammograms at 0, 45, and 100°C after I0

cycles at I00 mV/s. Nickel hydroxide oxidation and oxygen evolution are

within about I00 mV at 100°C compared with more than 200 mV at 0°C. The

separation decreases primarily because of the shift in the equilibrium

potential for oxygen evolution to more negative values and a decrease in the

overpotential for this reaction. Furthermore, at higher temperatures, the

anodic peak shifts closer to the oxygen evolution line with increasing cycle

time (Figure 5). In a relatively short time, the anodic peak disappears

completely from the voltammogram trace (Figure 6), although the presence of

217
the cathodic peak indicates that nickel hydroxide oxidation occurs

simultaneously with oxygen evolution. At 100°C, anodic peak shifting is even

more rapid, while at 0°C virtually no peak shifting with cycle time is

observed.

As noted by Macdonald and Owen (Ref. 21) and by McKubre and Macdonald

(Ref. 19), the reversible potentialrfor oxygen evolution is more negative than

that for Ni(OH)2/NiOOH at all temperatures of interest The appearance of the

nickel hydroxide . oxidation peak on the voltammograms is due to a high

overpotential for oxygen evolution. However, as the temperature is increased,

the overpotential is reduced, such that at 80°C and after extensive cycling a

distinct oxidation peak is no longer observed. This phenomenon may have

serious consequences for porous nickel electrode performance in concentrated

alkali solutions at elevated temperatures, because oxygen evolution in the

pores will occur simultaneously with charging. As noted previously, gas

formation within the pores may c0ntribute significantly to internal tensile

stresses that can rupture particle-particle ohmic contacts within the active

mass. Also, oxygen evolution represents a significant parasitic process that

will lower the coulombic efficiency of the porous electrode over a

charge/discharge cycle.

The anodic charge and peak current associated with nickel hydroxide

oxidation increase considerably at higher temperatures, suggesting that more

active material is present on the electrode surface. If film thickness is

assumed to be proportional to anodic charge, then thicker films are formed at

higher temperatures in a given number of cycles. A proportionality between

charge and film thickness is supported by the work of McKubre and Macdonald

(Ref. 19) where no evidence of film dissolution in rotating ring disc

experiments is reported. This indicates that an increased battery electrode

capacity might be anticipated if battery electrodes are operated at higher

temperatures, but of course any advantage may be offset by the decrease in the

oxygen evolution overpotential noted above.

Planar Nickel Electrodes Electrochemical Impedance Spectroscopy

The transmission line model requires a knowledge of the interfacial

electrolye/pore wall and electrolyte/backing plate impedances. It is assumed

in this work that these impedances can be described by the impedance of a

planar electrode in the same electrolyte. -This assumption can be supported by

several arguments. First, neglecting pore wall curvature, the basic structure

of the sintered battery electrode at the electrolyte/pore wall interface,

consisting ofthe me_al, film, and electrolyte, should be similar to that for

a planar electrode. The structure of the rolled and bonded porous electrodes

deviates somewhat from this geometry but it is similar if the graphite is

regarded as a substitute for the metal phase. Second, a concentrated

electrolyte was selected for this study, so electrolyte depletion within the

pores of the porous electrode should be minimal, particularly after one hour

at constant DC bias prio r to the measu[ement of the AC impedance spectra. The

electrolyte concentration at the pore wall should be approximately the same as

that at the film/electrolyte interface for a planar electrode.

218
In this study, an equivalent circuit for planar electro-oxidized (thin)

film electrode impedances is used as the interfacial impedance input to the

transmission line model. It can be argued that planar thick film electrode

impedances should be used, since the active material in nickel battery

electrodes is typically chemically or electrochemically deposited to a

relatively large thickness within the pores. However, thick films themselves

can be porous. Electro-oxidized thin films have comparatively smooth surfaces

and are more suitable for use as interfacial impedances in the transmission

line model.

Impedance spectra were recorded in sequences of increasing applied DC

bias. Figures 7 and 8 show a typical sequence of ambient temperature Bode

plots of log IZ[ vs. log _, and phase angle vs. log _, respectively, where IZI

is the impedance magnitude, _ is the angular frequency (2_f), and the phase

angle is the arc tangent of the ratio of the imaginary and real parts of the

measured impedance. Impedance magnitudes are found to decrease with

increasing potential , and a large decrease is observed when the nickel

oxyhydroxide phase is formed at 500 mV. Phase angles generally show two

maxima within the frequency range studied. The high frequency maximum shifts

sharply to lower frequencies at 500 mV corresponding to the film

transformation from nickel hydroxide to oxyhydroxide. However, both the

magnitudes and the phase angles exhibit similar features above and below the

nickel hydroxide/oxyhydroxide transition.

Impedance data were relatively unaffected by the potential step

increments used in the DC bias sequences. In one test sequence, an electrode

was polarized at 0 mV for one hour prior to an impedance spectrum measurement,

then was cycled potentiodynamically between -800 mV and 600 mV at i00 mV/sec

for 90 minutes. Following this, the impedance spectrum was again recorded at

0 mV after a one hour polarization at this potential. Both impedance spectra

were virtually identical indicating that prior polarization to higher

potentials does not significantly affect the planar electrode impedance

spectra. It also suggests that charge/discharge cycling in of itself will not

significantly affect the solution/pore wall impedance in the transmission line

model.

Planar electrode impedance spectra were also recorded at other

temperatures (Figures 9 and i0). Spectra were similar at all temperatures

from 23°C to 100°C. However, at O°C the high frequency relaxation shifted

markedly to lower frequencies, and only the beginning of the low frequency

relaxation is observable at the minimum frequency employed (6 mHz). At higher

temperatures, higher minimum frequencies were used to avoid data scattering

from noise, and again only the beginning of the lower frequency relaxation is

observable. Despite the restricted view of the low frequency relaxation, the

features of the elevated temperature data appear very similar to the ambient

temperature spectra discussed above.

The planar electrode impedance spectra described above were modeled with

the equivalent circuit shown in Figure ii. Mathematical impedance expressions

derived from this circuit were used in the transmission line model, together

with best fit component values (ie. capacitances, resistances and Warburg
coefficients) obtained from a curve fitting technique.

An (infinite thickness) Warburg diffusion impedancewas used in the

equivalent circuit because previous results by Chao, Lin, and Macdonald (Ref.

22), Liang et al (Ref. 23), Madouand McKubre(Ref. 8), and Zimmermanet ai.

(Ref. 7) indicated that a diffusion impedancedominated the low frequency

spectra over a wide potential range. Best fit componentimpedancevalues in

each of the equivalent circuits were determined by minimizing the weighted sum

of squares differences between the experimental and calculated impedance

data. An example of fitted data using the circuit shownin Figure ii and a

spectrum recorded at 0 mVDCbias is shownin Figure 12. Clearly the

essential features of the experimental data are reproduced by the model.

Table 1 lists the best fit equivalent circuit parameter values at several

potentials.

Porous Rolled and Bonded Eiectrodes

The transmission line model (TLM) contains eight independent variables.

Two of these variables are the pore wall and backing plate interfacial

impedances that are taken as planar nickel electrode impedances in this

study. Planar electrode impedances are each described by four component

impedance elements that may be frequency and DC bias dependent, as described

in the previous section. A total of 16 independent variables are used to

describe the impedance of porous nickel battery electrodes.

The TLM is used in this section to model porous electrode degradation

processes. Each of the 16 variables in the model were first determined as

accurately as possible in separate experiments. Impedance spectra were then

calculated from the model and were compared to porous nickel electrode

impedance spehtra. Adjustments to key parameters allow the calculated spectra

to progressively follow experimental impedances during degradation induced by

galvanostatic cycling at various temperatures. The TLM clearly showed cycle

dependent trends in several variables, and these trends are consistent with

experimental observations, as described below.

Initial estimates of transmission line model parameters for rolled and

bonded electrodes were selected as follows:

l. Pore wall and backing plate impedances, Z and Z', can be represented

by planar electrode frequency dispersions-fr6m either experimental

data or best fit calculated data derived from the equivalent

circuits shown in Figure ii. The circuit shown in Figure Ii and its

best fit component values (Table i) at 0 mV DC bias were selected

for this work.

. Scanning electron microscope examination of rolled and bonded

electrodes indicated that abowt 15000 pores were visible (at 100X)

on electrodes of area 1.27 cmZ given two ,,conditioning,,

charge/discharge cycles to their rated capacity. The smallest pores

on the electrode surface could not readily be resolved at IOOX and

were not counted.

220

=

Z

=

Z

Z

Z

=
.

4.

The initial average pore length, 2, was approximated as the

thickness of the electrodes (0.Ii cm).

The projected electrode area was 1.27 cm 2 for rolled and bonded

electrodes.

5° The total surface area of a rolled and bonded electrode after two

conditioning cycles was found to comprise approximately 33% pores.

. The KOH concentration in the pores was assumed to be constant along

the length and radius of the pores since a high KOH concentration

was used. While this assumption may not be strictly obeyed, the

KOH/H20 conductivity data of Lown and Thirsk (Ref. 22) shows that a

60% increase in conductivity is realized by halving the KOH

concentration from 8 molar to 4 molar. This is within a factor of

two. Therefore, the solution resistivity was approximated as a

constant of 3°i ohm cm for an 8 molar solution from the Lown and

Thirsk data.

. The solid phase resistivity, Pm' was difficult to estimate. The

rolled and bonded electrodes contain principally graphite and nickel

hydroxide. Graphite has a resistivity of about 0.0014 ohm cm (Ref.

25) while the resistivity of nickel hydroxide can be as much as ten

or twelve orders of magnitude larger at 0 mV as indicated by planar

electrode impedances. In this work, the graphite was assumed to

determine the solid phase resistivity, so an initial estimate of Pm

was taken as I0 ohm cm.

Bode plots calculated by direct substitution of the above

parameters in the TLM were compared with impedances for a rolled and

bonded electrode cycled twice at 23°C, with 40 minutes charging to rated

capacity, and 20 minutes for 100% depth of discharge (DOD). The

experimental impedance spectrum was measured after the electrode was

held in the fully discharged condition at 0 mV for one hour. The

features of the experimental impedance spectrum were present in the

calculated spectrum. However, the fit of the calculated curve was

relatively poor. Parameter value assignments were adjusted with a

design optimization software package (OPTDES) written at Brigham Young

University. OPTDES uses a set of design variables to minimize or

maximize one or more objective functions defined by the user in a usersupplied Fortran subroutine. Four optimization algorithms were used in

succession for each curve fitting.

OPTDES was used to minimize the sum of squares residuals (in the

Nyquist plane) defined as the objective function in the Fortran

subroutine. In a typical optimization sequence, interfacial and backing

plate impedances for 0 mV applied DC bias plus the electrode area were

first held constant, while the remaining variables were optimized. Then

some or all of the component impedances in the interfacial and backing

plate impedance expressions were allowed to vary along with the other

221
design variables.

Curve fitting procedures indicated that small adjustments to

several parameters significantly reduced the sum of squares error

between the experimental and calculated impedance spectra (Figures 13

and 14). These parameter adjustments from the initial value assignments

are discussed below.

I. The pore length, _, required for the "best" fit of the

experimental data is about three times the electrode thickness

(0.35 cm compared with 0.II cm thickness). This suggests that

tortuosity along the pores increases the active pore length by

this amount.

.

.

4,

The best fit number of pores is about 23000, whereas about

15000 pores could be resoived in a iOOX SEM photomicrograph.

This may indicate that smaller pores (not counted in the

photomicrograph) are not inactivated by their relatively

larger solution resistance, Psi/a, where a represents the

cross sectional area of a pore.

The optimized surface coverage of pores (i-O) was 0.33.

is in good agreement with the initial estimate of 0.3.

This

From Equation 4, the average pore radius can be calculated

from the optimized number of pores and the surface coverage

fraction of pores given above. The calculated average radius

is 24 #m, which agrees well with the pore sizes observed in

the SEM photomicrographs.

.

.

The optimized resistivities of the solid and solution phases

were 9.3 and 4.5 ohm cm, respectively. Predicted values were

I0 and 3.1 ohm cm, again in reasonable agreement.

The pore wall impedance, Z, required a small change to improve

the fit with experimental data. Specifically, the smaller

(high frequency) capacitance was decreased to about 7 _F from

43 _F at 0 mV DC bias. The latter number was obtained from

the 0 mV, 23°C planar electrode impedance spectra as described

in tee previous section. Smaller capacitances (around 29 #F)

were Obtained in a similar manner for planar electrode

impedances at lower potentials (-150 mV). Therefore, the

smaller capacitnce required to fit the porous electrode data

at 0 mV may be due to the potential drop across the porous

electrode. The pore wall impedance input to the TLM

represents an interfacial impedance averaged along the pore

wall. If a potential drop occurs along the pore, then the

pore wall impedance expression is selected at the average

potential. The decreased (optimized) capacitance suggests

that the average potential is lower than 0 mV, and therefore,

that a potential drop exists across the electrode.

222

i

z
Experimentally, the porous electrode is polarized anodically

at 0 mV, so the sign of the potential drop predicted by the

TLM is in agreement with that imposed on the electrode in the

cell.

7. The backing plate impedance, Z', had little effect on the

shape of the calculated impedance spectra. This result

supports the previous findings of a large pore length, and a

large number of pores. Both of these parameters are

proportional to the average solution resistance per pore

(Psn2/(l-e)A) and when large cause a redirection of current to

the solid phase (away from the solution/backing plate

interface.)

. The remaining variables in the pore wall interfacial impedance

(specifically the resistance, R, capacitance, CI, and Warburg

coefficient in Figure 12 were not optimized. Attempts to

optimize these variables always resulted in slightly smaller

sum of squares error, but with simultaneous distortions of the

calculated spectrum. The optimization software found pathways

to reduce the sum of squares error by skewing the features of

the calculated spectrum. This may indicate that the best fit

curves shown in Figures 13 and 14 represent only localized sum

of squares minima and that a better fit might be found. More

likely, it shows only that the optimization software can, in

some cases, reduce sum of squares residuals on curves with

complex shapes by unrealistic (albeit creative) manipulation

of a large number of variables. In this study, the

reproduction of the essential features of the experimental

impedance spectrum is considered more important than a smaller

sum of squares error.

Cycle Dependence

Impedance spectra for rolled and bonded electrodes at ambient

temperature are shown in Figures 15 and 16 after 2, 12, and 27

galvanostatic charge/discharge cycles. Impedance magnitudes increase

with cycle number at intermediate frequencies, while phase angle maxima

decrease at low frequencies and increase at high frequencies. This

behavior can be modeled by optimizing TLM parameters to minimize the sum

of squares error between experimental and calculated impedance data.

Parameter adjustments are discussed below:

. The optimized active pore length decreased from 0.35 to 0.28

to 0.20 cm after cycles 2, 12, and 27, respectively. A

considerable amount of the active material had spalled from

the electrode by the end of 27 cycles, and was found scattered

throughout the test cell. This could account for the

decreased pore length. However, the electrode also swelled

during the test, and after 27 cycles, the net thickness at the

center of the electrode was actually larger than the original

223
o

.

o

.

.

thickness. The active material remaining on the surface of

the electrode was easily flaked off, suggesting that the solid

phase particles farthest from the backing plate were not

active. The solid phase resistivity between these outer

particles may have been so large that they did not participate

in the electrochemical processes, and the pore length

effectively decreased.

Cycle dependent adjustments to the solid phase resistivity,

Pm' supports the above assertion that the outer particles were

not active. The average solid phase resistivity changed from

8.9 to 13.2 to 12.4 ohm cm after 2, 12, and 27 cycles. The

increase from the initial value is readily explained by

particle-particle bond damage or breakage during the swelling

process. Sloughing of the outer particles suggests that a

resistivity gradient developed across the electrode, with the

highest resistivity at the outer particles.

Optimized solution resistivities, Ps, remained virtually

constant throughout the cycling process. Resistivities after

2, 12, and 27 cycles were 4.5, 4.9, and 4.0 ohm cm,

respectively. Cycle independence was predicted since the KOH

concentration was large making electrolyte depletion or

exhaustion within the pores unlikely.

TLM modeling showed that the number of active pores, n,

decreased from 23000 to 8500 to 3300 during this cycling

sequence. A reduced number of active pores with cycling might

be explained by "restructuring" and solid phase particleparticle bond breakage which inactivates many of the pores.

The optimized pore coverage fraction, (1-8), decreased

slightly during cycling, from 0.33 to 0.27, and 0.25 after 2,

12, and 27 cycles. This is also consistent with the gradual

inactivation of pores during cycling.

According to the TLM, the potential drop across the porous

electrode diminished somewhat with increasing cycle number.

The optimized smaller capacitance in the pore wall interfacial

impedance increased from 7 uF after cycle 2, to 34 uF after

cycle 27. This indication of a cycle dependent decreasing

potential drop is also readily explained. The resistance of

the solid and solution phases per pore (Pmnl/eA and #_ n2/(le)A, respectively) can be regarded as indicators of the

potential drop across the porous electrode at a given DC

bias. Both resistances decrease with increasing cycle number

principally because of the decreasing pore length and number

of pores, and the potential drop decreases accordingly.

224
Temperature Dependence

AC impedance spectra for rolled and bonded electrodes at 0 mV were

measured at 0 and 40°C during their cycle life. The data had the same

features as those at ambient temperature, suggesting that electrode

degradation processes are similar at these temperatures.

No rolled and bonded impedance spectra were successfully recorded

at 60°C or at 100°C. Electrodes failed in less than one or two cycles,

and each test was terminated before impedance spectra were recorded.

Failure was associated with severe spalling of the active material and

the inability of the electrode to carry the imposed galvanostatic

current within the output voltage range provided by the galvanostat.

The cyclic voltammetric results presented earlier show that oxygen

evolution occurs rapidly during nickel hydroxide oxidation at 60°C and

higher temperatures. Oxygen evolution and the associated bubble

pressures within the pores are apparently major contributors to the

rapid electrode degradation at higher temperatures. Slow swelling and

sloughing at lower temperatures may also be related to the relatively

slower oxygen evolution reaction rate.

Porous Sintered Electrodes

The effect of cycling current, cycle number, and temperature on the

impedance data for a sintered electrode are shown in Figures 17-20. In

all cases, the impedance components vary in a relatively simple manner

with log _. Log |Z I typically decreases approximately linearly and then

becomes constant with increasing frequency, while 4 exhibits just one

maximum over the frequency range studied. The decrease in 4 to

approximately 0 ° and the corresponding independence of IZI with

frequency indicates that the impedance of the electrode is largely

resistive for a large part of the frequency range.

Effect of Cycle Number

The impedance spectra were essentially independent of cycle number

over the temperature range from 0 ° to 60°C, as shown in Figures 17-19.

(The small shift in the impedance at O°C is believed to be due to oxygen

contamination late in the electrode cycle life.) Even up to failure,

the spectra over this temperature range showed only small changes with

cycling. The only indication of an impending failure was a slight

fluctuation in the electrode potential just preceding failure.

Moreover, the sintered electrodes did not exhibit any sloughing or

swelling after failure as did the rolled and bonded electrodes.

Effect of Charge/Discharge Current

Impedance spectra obtained after cycling at high currents (26-min

charge to capacity and 14-min for 100% DOD) and low currents (80-min

charge to capacity and 40-min for 100% DOD) are shown in Figure 20. The

225
frequency response of IZI and _ were virtually identical for the two

sets of cycling currents, indicating that there was no discernible

degradation of the electrode structure within the range of currents

used.

Effect of Temperature

The shapes of the Bode plots were independent of temperature, but

the values of the impedance components did show some variation. The

principal changes were (i) a decrease in the value of log IZI at low

frequencies, (2) an increase in the frequency of the phase angle

maximum, and (3) a shift in the inflection of the log IZI curve to

higher frequencies.

At IO0°C, failures occurred in the first few cycles even at low

cycling currents based on an 80-min charging time. The impedance data

at this temperature exhibited scatter due to noise, and measurements of

the second harmonic indicated that the electrode response to a I0 mV rms

AC perturbation was not linear. The AC perturbation signal was reduced

to 3 mV, but the second harmonic was still too large.

The impedance spectra for the sintered electrodes differed from

those for the rolled and bonded electrodes in several respects.

Galvanostatic cycling of a rolled and bonded electrode, unlike a

sintered electrode, has a marked effect on the impedance components, as

shown in Figures 15 and 16. In addition, the curve for log |Z I shows

more than one inflection, and 4 exhibits two maxima over the frequency

range studied. The first maximum in 4 occurs in a similar frequency

range to that for the sintered electrode, implying that it is associated

with the active material whereas the second maximum is evidently related

to the electrode structure.

The data for the rolled and bonded electrode, as discussed above,

can be represented by a transmission line model, and parameters such as

pore length and solid phase resistivity obtained from the model changed

with cycle number. However, the impedance data for the sintered

electrodes were not analyzed using the TLM because the changes in the

impedance spectra were not sufficiently clear to render useful

information on the degradation mechanism. The sintered electrodes are

clearly more robust than are the rolled and bonded electrodes and hence

are much more resistant to restructuring and rupture of the ohmic

contacts in the active mass. Accordingly, we do not believe that our

failure to detect changes in the impedance spectra represents a

shortcoming of electrochemical impedance spectroscopy for studying

electrode degradation phenomena but simply that the sintered electrodes

are resistant to restructuring under the cycling regime employed in this

work.

226
SUMMARYANDCONCLUSIONS

Transmission line modeling results indicate a set of parameter

changes with cycle number that are consistent with experimental

observations. Initial estimates and measurementsof individual

parameter values comparefavorably with parameter values determined by

curve fitting to real rolled and bonded electrode impedancesat 0 mVand

at 23°C. Rolled and bonded electrode impedancesmeasuredat 0 and 40°C

behave similarly to those at ambient temperature, and similar parameter

changes with cycle number are indicated.

Specific changeswith cycle numberat 23Care:

i. The average pore length decreases with cycle number, but

always remains larger than the thickness of the electrode.

° The average solid phase resistivity increases with cycle

number.

. The solution resistivity within the pores remains virtually

unchanged during the cycle life of the electrodes.

4. The number of active pores decreases during cycling.

° The average resistance per pore of the solution and solid

phases decreases during cycling, and the potential drop across

the electrode decreases accordingly.

. The total porous electrode impedance is relativley insensitive

to the solution/backing plate interfacial impedance. This

indicates that little current flows along the entire pore

length in the solution.

Rolled and bonded electrodes break down rapidly when cycled at 60

and 100°C. Cyclic voltammetric results at elevated temperatures show

that the oxygen evolution reaction proceeds at a significant rate

concurrent with the electrode charging reaction after the first few

voltammetric cycles. Rapid rolled and bonded electrode breakdown during

galvanostatic cycling at elevated temperatures is probably due to

parasitic oxygen evolution processes.

Conversely, sintered electrode impedances do not change during

galvanostatic cycling, and failures occur abruptly after a relatively

large number of cycles. However, breakdown also occurred rapidly at

100°C, indicating that oxygen evolution processes may also affect

sintered electrode cycle life.

227
ACKNOWLEDGMENTS

The authors gratefully acknowledge the financial support of this

work by The U.So Dept. of Energy through LBL under Contract No. 712955.

REFERENCES

,

°

°

4.

5

6.

°

.

.

I0.

Ii.

Katz, M. H., McLarnon, F. R., Cairns, E. J., in Battery Division

Extended Abstracts, The Electrochem. Sot., Fall Meeting, Oct. 17-

21, 1982, p. 24.

Silverman, D. C., Corrosion , Vol. 37, No. 9, Sept. 1981, pp. 546-

548.

Glarum, S. H., Marshall, J. H., J. Eleotrochem. Soc., Vol. 129.

No. 3, March 1982, pp. 535-541.

r

Weininger, J. L., Breiter, M. W., J. Electrochem. Sot., Vol. iii,

No. 6,June 1964, pp. 707-712.

Sathyanarayana, S., Venugopalan, S., Gopikanth, M. L., J. Applied

Electrochemistry, Vol. 9, 1979, pp. 125-139.

Zimmerman, A. H., Janecki, M. G., in "Proceedings of the Symposium

on The Nickel Electrode," Gunther, R. G., and Gross, S. eds., JES,

Vol. 82-4, 1982, pp. 199-215.

Zimmerman, A. H., et al., J. Electrochem. Soc., Vol. 129, No. 2,

Feb_ 1982, pp. 289-293.

Madou, M. J., McKubre, M.CoH., J. Electrochem. Soc., Vol. 130, No.

5, May 1983, pp. 1056-1061.

De Levie, R., Electrochimica Acta, Vol. 8, 1963, pp. 751-780.

De Levie, R., in "Advances in Electrochem. and Electrochemo Eng.,"

Po Delahy, ed., Vol. 6, Interscience, New York, 1969, pp. 329-398.

eenhart' _SI J., Chao, C° Y., Macdonald, Do D., in "Proceedings

16th Intersociety Energy Conversion Engineering Conference, ASME,

New York, Aug. 9-14, 1981, pp. 663-666.

Park, J. R., Macdonald, D. Do, Corrosion Science, Vol. 12, No. 4,

1983, pp. 295-315.

Tuomi, D., J. Electrochem. Soc., Voi. 112, No. i, Jan. 1965, pp.

1-12.

228
14.

15.

16

17

18

19

20

21

22.

23.

24.

25.

Weininger, J. L., in "Proceedings of the Symposium on the Nickel

Electrode," Gunther, R. G. and Gross, S. eds., J. Electrochem.

Soc., Vol. 82-4, 1982, pp. 1-18.

Takehara, Z., Kato, M., Yoshizawa, S., Electrochimica Acta, Vol.

16, 1971, pp. 833-843.

Maskalick, N. J., Buzzelli, E. S., in Battery Division Abstracts,

The Electrochem. Soc., Fall Meeting, Oct. 17-21, 1982, p. 24.

Pound B. G., Singh, R. P., Macdonald, D. D., J. Power Sources,

Vol. 18, p. 1 1986.

Pickett, D. F., U.S. Air Force Technical Report AFAPL-TR-75-34,

Part I, Oct. 1975.

McKubre, M.C.H., Macdonald, D. D., J. Energy, Vol. 5, 1981, p.

368.

Sattar, M. A., Conway, B. E., Electrochemical Acta, vol. 14, 1969,

pp. 695-710.

Macdonald D. D., Owen D., "The Electrochemistry of Nickel Metal in

Lithium Hydroxide Solutions at 22, 170, and 250°C. '' High

Temperature High Pressure Electrocemistry in Aqueous Solutions,

Eds. R. W. Staehle, D. deG. Jones, and J. E. Slater, National

Association of Corrosion Engineers, Houston, 1976.

Chao, C-Y, Lin, L.F., Macdonald, D. D., J. Electrochem. Soc., Vol.

129, 1982, 1874.

Liang, R., Pound, B., Macdonald, D. D., Electrochem. Soc., Ext.

Abstr., 84-2, 1984, 356

Lown O. A., Thirsk H. R., Trans. Farad. Soc., Vol. 67. No. i,

1971, p. 132.

Handbook of Chemistry and Physics, Weast R. C. ed., CRC Press

Inc., 1974, p. F 159.

229
,.(m

E..w

f_

0

E-, ,_

u

f-wI._c

(,,,,I

:;,, ,,

_e,J

_rj

_mQ

N

z:t_

{..,)

mm

:=.c_

v

{,..).,_

mm

O

O

=0

O

O

..,I"

o O

O

O

O

N

o

O O

i •

r... =O

O

N

O

O

O

8

v

O O O O O

230
/ ///////_/

%%=. SOLUTION

N CHARGESTORAGE,MPEDANCE Z __

ELECTRODE

BACKING PLATE

(current collectOr}

FIGURE 1 RIGHT CYLINDRICAL MODEL OF IDEAL SINGLE PORE

(ira)x, o

Iz,

RE I

gm* dem gm

( em)X=O /

Rm.dx Rm.dx Rm'dx / Rm.dx

di m im -di m

dl s 'l

is is -dis> >

Rs .dx Rs'dx

,,/ fz,

\

(is)x=o

Rs .dx Rs .dx

gm" dSm

Rm' dx ]._,

Rs. dx

es es" des x=£

X=O

g s * d8 s

I

A

v

M

FIGURE 2 DISCRETIZED FORM OF TRANSMISSION LINE MODEL FOR A POROUS BATTERY

ELECTRODE OF FINITE THICKNESS

em and es are potentials in the metal and solution phases respectively.

im and is are currents in the metal and solution phases, respectively.

I and I' are the total current and the current flowing across the electrode backing plate/solution

interface at the base of the pore, respectively.

RE and M designate the reference electrode and current collector locations respectively.

231
FIGURE 3

16

12

8

<

E

_- 4

z

LU

n"

nU 0

-4

I

w ._ 37 Minutes

61 Hours

19.1 Hours

t 1

-8 J ] I

- 1000 -500 0 500 1000

POTENTIAL vs Hg-HgO-BmKOH/mV

VOLTAMMOGRAMS FOR PLANAR Ni ELECTRODE IN 8m KOH + 1% LiOH AT 23°C

AS A FUNCTION OF CYCLING TIME. VOLTAGE SWEEP RATE = 100mV/s

2

t_

E

<

E

z

LU

rr

tr

0

-1

FIGURE 4

m.__ 0°C

45 ° C

100_C

/

I l 1 ,

-2

-1000 -500 0 500 1000

POTENTIAL vs Hg-HgO-SmKOH/mV

VOLTAMMOGRAMS FOR PLANAR Ni IN 8m KOH+ 1%LiOH AT VARIOUS

TEMPERATURES AFTER 10 CYCLES AT 100mV/s

232
6 I I I

FIGURE 5

*- 2

z

uJ

(E

0E

o

• _ 10 Cycles

20 Cycles

30 Cycles

!i

I I I

-500 0 500

POTENTIAL vs Hg-HgO-6mKOH/mV

-2

-1000 1000

VOLTAMMOGRAMS FOR PLANAR Ni IN 8m KOH+ 1% LiOH AT 80_C, 100mV/s

AND 10, 20 AND 30 CYCLES

Iz

uJ

rr 1

mr

r_

FIGURE 6

-1 -

-2

-800 -600

] I L I ] I t

I I I I I I

•-400 -200 0 200 400 600 BOO

POTENTIAL vs Hg-HgO-..SmKOH/mV

VOLTAMMOGRAM FOR PLANAR Ni in 8m KOH + 1% LiOH AT 80°C.

20 mV/s, CYCLE 81

233
FIGURE 7

r_ 3

E

Jr

_N •\ "% e\=

".\

"% \. ,_

\. \\

OmV

n D --

--o- 200 mV

--*- 400 mV

--'- 500 mV

\\\. "\\_%%

\•\ \%.

z,...r ,

=

-2 0 2 4 6

log _,_/radians

AMBIENT TEMPERATURE log IZ; vs Iog_ DATA FOR PLANAR Ni ELECTRODES

IN 8 m KOH-_ 1% LiOH OVER A RANGE OF APPLIED DC BIAS

8O

6O

40

Z

,<

I

ca.

2O

0

-2

FIGURE 8

--'- 400 mV °

500 mV

\

\

\

u

0 2 4 6

log i.._'/r ao, a rl$

Pt-IASE ANGLES CORRESPONDING TO DATA IN FIGURE 7

234
3

%E

1

-1

., -,- 23°C

\, • =\ --'- 60_C

"%"__i_:_,-.,oo°_

\\'\:\

\.\'\:.\

\._;_.\

\ =

\ +\ o"i.e

"_.\ \e "nn"=

i o'_% @'e- o

i i ,

-2 0 2 4 6

log _{ra_lans

FIGURE 9 IMPEDANCE SPECTRA AT VARIOUS TEMPERATURES FOR A PLANAR Ni

ELECTRODE IN 8m KOH + 1% Lj0H AT 200mV

100

8O

a;

== 60

uJ

,_I

(..3

Z

<

LUtn 40

"r

20

,,,¢1_ g--m-i= - g_

\ Z A-"/_'e_<".

.., ,i <'// \-\)

x?./_/" \V\

, XUj \.t'it

"%°/ " "'0.

_=_ --,,- O°C 23 ° C 't'Q'_'i!',Q=_ "

Z:Z _°c

100°C \

-2 0 2 4 6

log _-.:fradi_ln s

FIGURE 10 PHASE ANGLES CORRESPONDING TO DATA IN FIGURE 9

235
AAA

VVV

R=

C_ C2

R

FIGURE 11 AN EQUIVALENT CIRCUIT FOR PLANAR

NICKEL ELECTRODE IMPEDANCES

IN B m KOH + I% LiOH

r_ 3

E

t-

__c

N

_2

•,, / .,.-.-.\

)

o\

-•- Exp. '_\

--+- Ca]c. "\_,

i i

-2 0 2 4 6

log _._/radians

FIGURE 12 EXPERIMENTAL AND CALCULATED IMPEDANCE SPECTRA AT 0mV APPLIED

D_ BIAS USING THE EQUIVALENT CIRCUIT IN FIGURE 11 FOR IMPEDANCE

CALCULATIONS. (23°C)

=_

236
3

--+-- 0 mV, Catc.

--e-- 0 mV, Exp.

?

1

0 ! i i

-2 0 2 4

log ¢._/radians

FIGURE 13 EXPERIMENTAL AND CALCULATED log rZl vs Iogc_ DATA AT 0mY AND 23°C FOR

A ROLLED AND BONDED POROUS ELECTRODE CYCLED TWICE. (40rain CHG

TO CAPACITY AND 20rain TO 100% DOD.) CALCULATED SPECTRUM

DETERMINED FROM OPTIMIZED TLM PARAMETERS DESCRIBED IN TEXT.

8O

6O

UJ

.J

c_ 40

z

LU

:z:

20

/= = - o,,,, -e- 0 mY. Calc.

--o- 0 mY, Exp.

/ \.l

i i

-2 0 2 4 6

log _/radians

FIGURE 14 PHASE ANGLES CORRESPONDING TO DATA SHOWN IN FIGURE 13

237
E

t--

_m

N

___ 2 Cycles

--o- 11 Cycles

-*- 27 Cvcles

-_

-2 0 2 4 6

109 _,._/ritdians

FIGURE 15 EXPERIMENTAL log IZI vs Iog_ DATA FOR A ROLLED AND BONDED POROUS

ELECTRODE AT 0mV AND 23°C AFTER VARIOUS CYCLES. (40rain CHG TO

CAPACITY AND 20rain TO 100% DOD.)

8O

60

uJ

_J

(..o AO

z

¢[

I.U

¢n

c[

"r

20

FIGURE 16

_=_ 2 Cycles

m/'l\ -i- 12 Cycles

/ "_ -+- 27cyc,,s

•._:/_\; / \.

/ \\V/ _X

ti__trl-- II - l_

-2 0 2 4 6

tog l_/radlans

PHASE ANGLES CORRESPONDING TO DATA SHOWN IN FIGURE 17

238
I oÁ' I' I I I 90

4

3

N 45

13_

1

opI __ I 1 I ,,: -to

-1 0 1 2 3 4

bg (W)

1_-&5054.3-29

FIGURE 17 BODE PLOTS OF IMPEDANCE DATA FOR A SINTERED

ELECTRODE AT 0°C AND 0 V AFTER 5(11, • ) AND 186

( $, ,_ ) GALVANOSTATIC CYCLES

(26-rain. charge to capacity with 14-rain discharge

to 100% DOD).

239
4

3

N

_2

o

1 •

0

90

k

w

_J

45 z

<

LLI

<

I

n:

0

FIGURE 18

0 I 2 3 4

k_g(W)

RA-350543-30

BODE PLOTS OF IMPEDANCE DATA FOR A SINTERED

ELECTRODE AT 23°CAND 0V AFTER 5(II,• )AND 275

(0,• )GALVANO_STATICCYCLES

(26-min.chargetocapacitywith14-raindischarge

to100°/oDOD).

240
3

N

_2

o

0

-1

FIGURE 19

I I I I 1

0 1 2 3 4

log (W)

9O

45

u.I

_J

Z

uJ

cO

o.

t=_&-3,50543-31

BODE PLOTS OF IMPEDANCE DATA FOR A SINTERED

ELECTRODE AT 60°C AND 0 V AFTER 6( II, •) AND 172

( 0, A ) GALVANOSTATIC CYCLES

(26-min. charge to capacity with 14-min discharge

to 100% DOD).

24]
4 t I I I I I I 90

3

N 45 Z_

1

0 0

-1 0 i 2 3 4

FIGURE 20

RA-350543-32

BODE PLOTS OF IMPEDANCE DATA FOR A SIN'I-ERED

ELECTRODE AT 23°C AND 0 V AFTER CYCLING AT HIGH

CURRENTS(==, e) AND LOW CURRENTS( $, '=)

(High currents = 26-m_nl chargeto capacity with 14-rain

discharge to 100% DOD; low currents = 80-min charge

to capacity and 40-min discharge to 100% DOD).

242
KOH

N )O- 20475

CONCENTRATION EFFECT ON THE CYCLE LIFE OF NICKEL-HYDROGEN

CELLS. IV. RESULTS OF FAILURE ANALYSES

H.S. Lim and S.A. Verzwyvelt

Hughes Aircraft Company

Los Angeles, California 90009

KOH concentration effects on cycle life of a Ni/H 2cell have been studied by carrying out a cycle lif_. tp._.:

of ten Ni/H 2 boiler plate cells which contain electrolytes of various KOH concentrations. Failur_

analyses of these cells were carried out aftercompletion of the life test which accumulated upto 40.0¢_-,

cycles at an 80% depth of discharge overa period of 3.7 years. These failure analyses included studi_._,

on changes of electrical characteristics of test cells and component analyses after disassembly of thn

cell. The component analyses included visual inspections, dimensional changes, capacity measurements of nickel electrodes, scanning electron microscopy, BET surface area measurements, ar_

chemical analyses. Results have indicated that failure mode and change in the nickel electrode v_.ri,=_

as the concentration was varied, especially, when the concentration was changed from 31% or highp:-

to 26% or lower.

INTRODUCTION

Long cycle life has been demanded for a Ni/H 2cell at deep depth of discharge cycles used in ._vac_.

craft applications, especially for a low earth orbit (LEO). We have reported earlier that cycle life of I',]i/1-_

cells was improved tremendously by using lower KOH concentrations in electrolyte than convention,_'

31% KOH 1. The cycle life has been improved as much as nine times by simply changing the. cc, ncentr_

tion to 26% (by weight) from a conve ntional value of 31%. A cycle life test has been completed forten ,_l,

H2cells containing electrolytes of various KOH concentrations ranging from 21 to 36%. In the r;rp._-,;

paper, we report failure analysis results of these cells which include studies on chanG_.._ of c_-"

performance and cell components after cell disassembly. We have also studied electroch_rc,i; ;_

behavior of the cycled nickel electrodes. This result will be reported later.

EFFECTS OF KOH CONCENTRATION ON CYCLE LIFE

A cycle life test was carried out in a 45-min LEO regime at 80% depth of discharge on tP.n l'Ji/H_ boil, _

plate cells containing electrolytes of 21 to 36% KOH. Test period covered 3 years and 8 ,mo_th-:..=_

accumulated up to 39,573 life cycles. Details of the cycle regime have been reported earli_.r.' Lif_. to:.:

results are summarized in Table 1. All cells of relatively short cycle life (BP3 to 6) failed by Io,,, e_,d-._,'

discharge voltage (EODV), while all cells of long cycle life (BP1,2, 9, and 10) failed the cycling test d::

to a "soft" short formation which was indicated by rapid self-discharge. Fully charged ceils discF.a ;_.-!

completely in 30 to 35 h. The effect of the concentration on cycle life is shown in Figure 1 and 2. T =_.-

cycle life to 0.9 V of EODV peaked at 26% KOH showing tremendous improvement (about nine time_"

over the conventional electrolyte of 31% KOH, but the life decreased sharply when the concent,"atin:_

was lowered further below 26%. If the voltage criteria for failure was lowered to 0.5 V. however, KC.,I.i

concentrations as low as 21% also gave the extreme long life (Fig. 2). This rather drastic chn,n.qe '.'--

cycle life by a small change in the voltage criteria was due to development of a large caDacity h_'_,_ _

low discharge voltage (about 0.8 V) plateau which is often refered as "the second plateau. ''r243
CELL

NO.

BP1

BP2

BP3

BP4

BP5

BP6

BP7 °

BP8

BP9

BPI0

TABLE 1. LIFE TEST RESULTS OF NI/H2CELLSAT 80%

DEPTH-OF-DISCHARGE

[KOH]

%

21

26

26

31

31

36

21

26

26

23.5

NO.CYCLE NO.CYCLE TOTAL FAILURE

TO 0.9 V TO 0.5 V NO.CYCLES MODE

5,047

39,230

4,329

2,979

3,620

1,268

1,037

>30,549

23,598

4,803

>38,191

>39,573

9,241

3,275

4,230

1,845

6,508

>30,549

>24,594

28,495

38,191

39,573

9,241

3,286

4,230

1,845

9,402

30,549

24,594

28,495

"CYCLED AT 70% DEPTH-OF-DISCHARGE (DO0) FROM 1644 TO 4644

CYCLES AND AFTER 6508 CYCLES.

4OOO0

3OOOO

;>

m

o

_o

LU

-- 20000

._J

LU

o

SOFT SHORT

SOFT SHORT

LOW EODV

LOW EODV

LOW EODV

LOW EODV

LOW EODV

REMOVED

SOFT SHORT

SOFT SHORT

1O0OO

"Not Failed l

21 21 23.5- 26 26 26 3i 31 36

KOH CONCENTRATION, %

FIGURE 1. COMPARISON OF CYCLE LIFETO 0.9 V OF Ni/H2 CELLS WITH

VARIOUS KOH CONCENTRATIONS IN THE ELECTROLYTE.

-7

CHANGE OF CELL CHARACTERISTICS AFTER LIFE TEST

_, Series0f CiUpi_cate-t(_t_-pl_a-te Capacity measurements were carried out before and after the cycle

life test using various charge and discharge rates in order to evaluate the rate effects on cell capacity.

The charge rate effect was studied by measuring cell capacity at C/2 discharge rate to 1.0 V after

charging them at C/10 rate for 16 h, C/2 rate for 160 min, C rate for 80 min, and 20 rate for 40 mirl,

respectively. The discharge rate effect was studied by measuring cell capacity at various discharge

rates (C/10, C/2, 1.0C, 1.37C, 2.0C, 2.74C, and 4.0C) to 1.0 V aftercharging them at C rate for80 min.

Internal resistance of cells was evaluated from mid-discharge voltages at various discharge rates.

Charge Rate Effects on Cell Capacity - Initial and end-of-life (EOL) capacity values are plotted

against charge rates in Figure 3 and 4, The capacities of all cells improved asthe charge rate increased

from 0.1 to 0.5 C rate. This trend was more pronounced after the cycle test. Further increase in charge

244
0

t--

UJ

u,.

._1

LLI

o

4oooo

3OOOO

2OOOO

1oooo

* NOT FAILED I

23.5

###.

#-ziJ

#//_,

t!!_ t

FI## _,/#.,.I

f.F#-# flJ_

.... _///j

#//..

F_f /

/-!!_,

9"//, ....

tY/I/, ..... [

KOH CONCENTRATION, °N

21 31 31 36

FIG. 2. COMPARISON OF CYCLE LIFE TO 0.5 V OF Ni/H2 CELLS

WITH VARIOUS KOH CONCENTRATIONS IN THE ELECTROLYTE.

rate over0.5 C did not affect capacity of cells with high KOH concentrations (31 and 36%) while capacity

of cells with low KOH concentrations improved slightly at 1 C rate. Capacities of all cells leveled off for

the charge rates between 1 and 2 C. Although absolute values of the capacity decreased after life test

with exception of BP8 and 10, general pattern of the charge rate dependence did not change with the

cycle test.

Effects Of Discharge Rates On Capacity - Initial and EOL capacity values are plotted against

discharge rates in Figure 5 and 6. Capacity of cells in general decreased slightly as discharge rate

increased. This dependence ofthe capacity on discharge rate became more pronounced after the life

test. We had also observed such dependence eadier with another group of 31% KOH cells?

Cell Resistance- Plots of mid-discharge voltages (measured during the initial and EOL capacity tests)

against the discharge current gave good straight lines as shown in Figure 7 for an illustration. The

slopes of these lines represent the values of cell resistance. The cell resistances before and after the

cycle life test are summarized in Table 2. Cell resistance did not appear to have changed significantly

after the cycle life test except for BP1, BP3, and BP10. BP1 showed a significant reduction of the

resistance. BP3 which had a defective current collector tab on one of the H2 electrodes showed a

significant increase. Overall, the change in the cell resistance with cycling appear to be too small to

accou nt eitherforthe change of capacity dependence on discharge rate orto be the cause of cell failure.

COMPONENT ANALYSES AFTER CELL TEARDOWN

Afterthe post-life-test capacity measurements, test cells were charged for 80 min at C rate and then

discharged at C/2 rate to 0.5 V followed by shorting them with a 1.0-_ resistance for a minimum of 24 h

before disasse mbly of the cells forco mponent study. This study included visual inspections duri ng the

245
,&

r,

6.5 /" .... .....

0 ..........'"

13.0

5.5 O' ,, , -0-

4._

i I

0.0 1.0 2.0

CHARGE RATE, C

21% (BP1)

...... 40 ......... 26% (BP2)

......0 ..........26% (BP3)

----O ....... 31% (BP4)

31% (BP5)

36% (BP6)

21% (BP7)

....-.O ........ 26% (BPS)

...... O ........ 26% (BPg)

23.5% (BP10)

BP7 (EOL 9402-)

BP3 (EOL 9241~)

----e----- BP4 (3286-)

...... O ....... BP5 (4230~)

...... _ ........ BP6 (1845~)

FIGURE 3. PLOTS OF INITIAL AND EOL CAPACITIES VS. CHARGE RATES. CAPACITY WAS

MEASURED BY 0.5 C RATE DISCHARGE TO 1.0 V AFTER CHARGING AT VARIOUS RATES: 0.1 C

RATE FOR 16 H, 0.5 C FOR 160 MIN., 1.0 C FOR 80 MIN., OR 2.0 C FOR 40 MIN.

U

2

1 --

0

0.0

oJ o.............................. •.......................................................

- • _j"O

_0 BP1

........................... 0 BP2

O BP8

_ BP10

! _ i t I ! I f i I l ] t ] I ] f 1 I

0.5 1.0 1.5 2.0

CHARGERATE, C

FIGURE 4. PLOTS OF EOL CAPACITIES VS CHARGE RATES. CAPACITY WAS MEASURED

BY 0.5 C RATE DISCHARGE TO i.0 V AFTER CHARGING AT VARIOUS RATES: 0.1 C RATE

FOR 16 H, 0.5 C FOR 160 MIN., 1.0 C FOR 80 MIN., OR 2.0 C FOR 40 MIN.

246
-I8.0

6.0

5.5

5.0

4.5

4.0

75- ;

6.5

I

0.0 1.0 2.0 3.0 4.0

CHARGE RATE, C

21% (BP1)

......0 ........ 26% (BP2)

.... _._ 31% (BP4)

31% (BPS)

36% (BP6)

21% (BP7)

............ O .......... 26% (BP8)

......o ........ 26% (BPg)

23.5% (BP10)

-----4"----- BP7 (EOL 9402~)

BP3 (EOL 9241~)

BP4 (3286 ~)

........O .............BP5 (4230 ~)

....... O,........... BP6 (1845~)

I I

5.0

FIGURE 5. PLOTS OF INITIAL AND EOL CAPACITIES VS. CHARGE RATES. CAPACITY WAS

MEASURED BY DISCHARGING CELLS AT VARIOUS RATES TO 1.0 V AFTER 1.0 C RATE

CHARGE FOR 80 MIN.

6.0

8

o

t_

INITIAL

6 - O_ "d_ _-_"_'_-O

'qF

5 -- AFTER CYCLING

4 -

3-

2-

1 -

0 I I I I I I I I I I I I I I I I I I I I i

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0 3.2 3.4 3.6 3.8 4.0 4.2 4.4

DISCHARGE RATE, C

FIGURE 6. PLOTS OF INITIAL AND EOL CAPACITIES OF BP8 VS DISCHARGE RATES.

CAPACITY WAS MEASURED BY DISCHARGING THE CELL AT VARIOUS RATES TO 1.0 V

AFTER 1.0 C RATE CHARGE FOR 80 MIN.

247
1.4

LLI

(3

¢1:

O

CO

a

_.+

1.3

1.2

BPI

.... O .... BP2

O ..... BP8

--O'-'- BP10

1.1 t

0 t0 20 3O

DISCHARGE CURRENT, A

FIGURE 7. PLOTS OF MID-DISCHARGE VOLTAGES OF TEST CELLS

VS DISCHARGE RATES FOR END-OF-LIFE PERFORMANCE TEST.

TABLE 2. CELL RESISTANCE (R) MEASURED FROM MID-DISCHARGE VOLTAGES.

CELL KOH NUMBER INITIAL R EOL R CHANGE CELL P°

NO. % OF CYCLES m,'£ mQ mC-J PSlG

BP1

BP2

BP3

BP4

BP5

BP6

BP7

BP8

BPI0

2t .0

26.0

26.0

31.0

31.0

36.0

21.0

26.0

23.5

38,191

39,573

9,241

3,286

4,230

1,845

9,402

30,549

28,495

10.83

(5.8g)

7.53

7.15

7.09

7.61

7.30

6.59

6.57

6.36

6.26

9.69

8.26

6.96

8.49 •

7,14

6.71

8.24

- 4.47

+ 0.37

+ 2.16

+1+11

-0.13

+ 0.88

-0.16

+0.12

+ 1.67

615

705

130

150

225

160

130

0

635

• PRESSURE PRIOR TO DISASSEMBLY.

disassembly, electrode thickness measurements (without washing and drying), scanning electron

microscopy (SEM) of the cross-section of cell stacks, surface area measurements using BET techrLique, measurements of nickel electrode capacity and discharge voltages in flooded electrolyte cells,

_tnd chemical analyses of the nickel electrodes.

visual Inspecti0ns- Conditions of cell stack and its components have been inspected carefully during

........... disassem+bFy 07 Cycle_cells. These |nspections inCfudedWetness of c0mp0nents, +degree of loOse

ulack nickel active material ("black powder") on cell components, penetration of the active material

ti+_ruughthe separator, adhesion of hydrogen electrode to gas screen, and integrity of nickel and

I',yc_rogen electrodeS. Observati0n§which appear tobeirnportafit are summarized as follows:

Wetness of the electrode and the degree of"black powder" formation increased noticeably as one

3pproached the bottom of the stack (no.6 unit cell) from the top (no. 1 unit cell). The observation on

248
wetness change indicated that gravitational effect on the electrolyte distribution is substantial even

in this relatively short stack. The observation on "black powder" formation appears to indicate that a

wetter unit cell might have been used more heavily than a drier unit cell.

Electrodes from BP7 (21% KOH; 9402 cycles) showed noticeably less "black powder" on the

adjacent separator and in gas screen than those from high KOH concentration cells which had

lower number of cycles. The black powder in the gas screen was found only in the peripheral area of

the stack with BP7 while it was found all over the screen with the othercells which were cycled less

than BP7. However, the 21% KOH cell (BP1; 38,191 cycles) which were cycled much longer

showed heavy black powdering. It appears that rate of black powder formation is reduced as KOH

concentration decreased.

Electrical lead (tab) of #5 H2 electrode of BP3 was not firmly attached (defective weld) to the

electrode while its paired nickel electrode and separator appeared new (almost without the black

powder). This observation stronglyindicates that the cause of the poorperformance of this cell from

the beginning was due to this poor electrical contact. The BP3 should be considered as an

anomalous cell.

All nickel electrodes from cells which had extremely long cycle life (>24,594 cycles) showed

"mushy" appearance indicating that some active material has loosened and squeezed out of the

sinter structure. This indication was confirmed by scanning electron microscopic (SEM) pictures as

will be discussed laterin this report.

All cells which had extremely long cycle life (>24,594 cycles) showed the following: Outeredges of

gas screen were welded to Teflon layer of the H2electrodes. Damage with partially lost electrode

material were apparent on parts of peripheral area of electrode stacks. The outer coined nickel

electrodes was badly eroded. These observations indicated that violent reco mbi nations of O2with

H2had occurred around peripheral area of electrode stacks during the cycling test period.

Pressure Changes- Gas pressure of test cells in the fully discharged state were measu red just before

their disassembly using a mechanical pressure gauge. This was done in order to determine the net

pressure change during the cycle life test. Gas composition was analyzed using a gas chromatograph.

Gases in all cells were pure hydrogen within experimental error. Results on pressure readings are

shown in Table 2. Long cycled cells (BP1,2, and 10) showed roughly 470 psi higher pressure than

those of relatively short cycle life (BP4, 5, and 6). (BP8 might have leaked slowly.) Although the long

term sealing reliability of presently used boiler plate cell cases was questionable as discussed earlier, 'b

the pressure increase in long cycled cells has be real. This pressure increase appears to be due to

nickel corrosion from the sintered plaques probably according to equation (1):

Ni + 2H20 > Ni(OH)2 + H 2 (1)

Dimensional Changes of Stack Components - Thickness of nickel electrodes was measured by a

micrometer at three different locations on each electrode before and after cycling. The thickness after

cycling was measured both before and after rinsing out the electrolyte and drying. These two values

agreed with each otherwithin experimental error. The average values of these measurements foreach

electrode are tabulated in Table 3. Long cycled electrodes (BP2 and BP8) showed expansion values

249
upto 80% of the original thickness. These long cycled electrodes showed heavy extrusion of active

material out of the electrode intovoids of the gas screen. Typically imprints of gas screen areclearly

visiblewhen gas screen isseparatedfromthe electrode.The highvalue of expansion (upto 80%) was

due to contribution of this extruded material.

Average expansionvalues ofelectrodes from individualcells areplotted against numberofcycles in

Figure 8. The expansion values appear tofall on a straight line eventhoughthe individual cells have

different KOH concentrations. This observation appears to indicatethat the expansion rate on long

term cyclingis ratherinsensitiveto KOHconcentrationwhichiscontraryto ourearlierunderstanding.3,4,_

The expansion itselfdoes not appearto bethe maincause of electrodeor cell failure.

TABLE 3. NICKEL ELECTRODE EXPANSION AFTER CYCLE TEST IN NI/H 2 CELLS

WITH VARIOUS KOH CONCENTRATIONS.

CELL (%KOH) ELECTRODE THICKNESS AFTER THICKNESS BEFORE EXPANSION"

(CYCLE NO) I.D. CYCLING," MM CYCLING, MM %

BP1 (21%)

(31,191)

BP2 (26%)

(39, 573)

BP3 (26%)

(9641 )

BP4 (31%)

(3275)

BP5 (31%)

(4230)

BP6 (36%)

(1845)

BP7 (21%)

(9402)

BP8 (26%)

(30,549)

BP9(26%)

(24,594)

BP10

(23.5%)

(28,495)

28-0S(#1, TOP)

28-04(#2)

24-03(#5)

24-01 (#6, BOI'TOM)

20-01 (#1)

20-02(#2)

19.05(#5)

19-06(#6)

25-01 (#1)

26-02

35-03(AN)

35-04(#6)

25-o9(#1)

26-03

26-09

27-07 (#6)

22-02 (#1)

22-03

22-09

20-03 (#6)

27-03 (# 1)

27-05

35-08

35-09(#6)

05-05(#1)

06-06(#2)

21-07(#3)

21-08(#4)

21-o5(#1)

21-01(#2)

03-06(#5)

03-02(#6)

03-03(#1)

04-07(#2)

04-07(#5)

06-07(#6)

04-03(#1)

16-05(#2)

16-02(#5)

16-09(#6)

1.166

1.141

1.268

1.225

1.310

1.274

1.341

1.327

0.930

0.933

0.762

0.947

0.837

0.814

0.819

0.778

0.809

0.814

0.785

0.758

0.816

0.777

0.785

0.771

0.876

0.916

0.929

0.949

1.225

1286

1.297

1.304

1.217

1.229

1.233

1.208

1.330

1.281

1.353

1.285

0.761

0.758

0.737

0.751

0.732

0.739

0.737

0.735

0.768

0.744

0.750

0.753

0.727

0.732

0.755

0.744

0.738

0.732

0.721

0.755

0.744

0.720

0.724

0.723

0.748

0.746

0.755

0.755

0.728

0.762

0.727

0.719

0,736

0.735

0.741

0.752

0.751

0.736

0.760

0.746

MEASURED VALUES INCLUDE EXTRUDED ACTIVE MATERIAL INTO VOIDS OF GAS SCREEN.

53.4

50.5

71.8

J63.2

79.0

72.5

81.8

76.1

21.1

25.3

1.6

25.7

15.1

11.2

8.5

4.6

9.6

13.6

8.9

0.4

9.6

7.8

8.4

6.7

17.2

22.7

23.1

25.6

58.2

68.7

78.4

81.3

65.4

63,2

66.4

60.6

77.2

74.1

78.0

72.1

250

=

_Z=
Thickness of stack components including separator and gas screen was also measured from SEM

picture of cross-section of cell stack. The results are shown in Figure 9 as a function of the number of

cycles. Dimension of unit cell was generally unchanged with cycling. Nickel electrode expansion was

again roughly linearwith the number of cycles. Separator and gas screen were compressed as the

nickel electrode expanded.

8O

[] []

[]

20 []

0

0

FIGURE 8.

, I , I ,

10000 20000 30000 40000

NUMBER OF CYCLES

A PLOT OF NICKEL ELECTRODE EXPANSION VS. NUMBER OF CYCLES.

2.4

2.2

2.0

1.8

1.6

1.4

1.2

1.0

0.8

0.6

0.4

0.2

0.0

UNff CELL

I I I

0 1(3000 20000 30000 40000

NUMBER OF CYCLES

FIGURE 9. PLOTS OF CELL STACK COMPONENTS

THICKNESS AS A FUNCTION OF CYCLING.

251
TABLE 4. FLOODED CAPACITIES OF NICKEL ELECTRODES FROM CYCLED CELLS. #1 : TOP

ELECTRODE OF THE STACK. #6 BOTTOM ELECTRODE OF THE STACK.

A NEW ELECTRODE HAD 1.24 AH FOR AN AVERAGE VALUE.

1.17

1,18

0.99

1.16

0.98

0.98

0,95

121

1.26

KOH

CONC.

CELL

NO,

%

BP1 21

BP2 26

BP3 26

BP4 31

BP5 31

BP6 36

BP7 21

BP8 26

BP10 23,5

CAPACITY OF #1 ELECTRODE. AH

!

1ST I 2ND

1.18

1.17

1.00

1.11

0.96

1,08

0.93

1.15

1,26

"iAFTER CHANGING ELECTROLYTE)

CAPACITY OF #6 ELECTRODE, AH

BP1

BP2

BP3

BP4

BP5

BP6

BP7

BP8

BP10

3RD AV, 1ST 2ND 3RD AV.

1.18 1.24 1.24 1,24

1.17 1.25 1.24 1.25

1.00 1.00 0.70 0,69 0.70 0.70

1.12 1.13 0.73 0.77 0.79 0,76

0.98 0.97 0.69 0.70 0.71 0.70

1.17 1.08 0.69 0.81 0.81 0.77

0.93 0.94 1.01 1,02 1.01 1.01

1.18 1.13 1.11 1.12

1.26 1.24 1.24 1.24

CAPACrI-_

RATIO

#61#1

1.05

1.07

0.70

0.67

0,72

0,72

1.08

0.95

0.98

31 1.11 1.10 1.1 t 1.16 1.15 1.16 1,04

31 1.13 1.23 1.18 1.16 1.15 1.16 0.98

31 1.02 1.04 1.03 1.03 0,75 0.77 0.77 0.76 0.74

31 1.14 1.14 1,14 1.14 0,79 0.80 0,80 0.80 0.70

31 1.01 1.03 1.03 1,02 0,72 0.72 0,72 0.72 0.72

31 1.04 1.04 1.03 1,04 0,69 0.70 0,70 0.70 0,68

31 1,02 1.06 1.05 1.04 1.03 1.0g 1,08 1.07 1,03

31 1.21 1.20 1.21 1,06 1.05 1.06 0.88

31 1.21 120 1.21 1.21 1.20 1,21 1.00

Electrode Capacity in a Flooded Cell - Capacities of two nickel electrodes from each failed cell were

measured in a flooded cell at C/2 discharge rate after charging C/10 for 18 to 19 hours. One electrode

was from the top of the cell stack (#1) and the other was from the bottom of the stack (#6). Two to three

capacities were measured in an electrolyte of the same KOH concentration as the original test cell and

then all electrolytes were changed with 31% KOH for additional three capacities. The results are

summarized in Table 4. With and exception of BP3 which waSananomalous ceil, the capacity loss of

nickel electrodes in low KOH concentration (21 to 26%) cells was relatively low and independent of the

electrode position (#1 or #6) despite long cycling. With higher KOH concentrations (31 to 36%),

however, the capacity of bottom (#6) nickel electrode was roughly 30%smallerthan top (#1) electrode

of the same cell. It appears that the cause of this faster degradation of the electrodes from the bottom

(#6) than those from the top (#1) of the stack was originated from the variation of the electrolyte content

from the top to the bottom. Cell teardown showed that the bottom part (#6) of the stack had much more

electrolyte than the toppart (#1) due to gravitational segregation. The wetter electrode pair (#6) is

expected to pass higher current throughout the cycle test than the drier pair (#1) which was in parallel

electrical connection. This uneven current distribution would have caused heavier usage of the #6

electrode than the #1. This heavy usage appears to cause a severe capacity decrease in higher KOH

concentrations while the decrease is relatively minor in 10WKOH concentrations. No apparent relationship was fou nd between the capacity decrease and the electrode expansion.

Chemical Analyses- Chemical Analyses resultsof electrolyte composition of cycled cells are shown

in Table 5. The overall electrolyte concentrations in long cycled ceils (BP2, BP8, and BP10) was increased by 3 to 4% from the initial value, indicating possible nickel plaque corrosion. Carbonate concentration which was negligible initially increased slightly in all cycled cells. The amount of carbonate

build-up was roughly proportional to logarithmic number of cycles as shown in Figure 10. A possible

source of carbonate might be corrosion of plaque which may contain about 0.1% (by wt.) of carbon because it is made of carbonyl nickel powder (INCO Type 287).'

252

z

=
TABLE 5. COMPOSITION OF ELECTROLYTE FROM CYCLED CELLS.

INITIAL CONC. KOH, K_CO3, TOTAL K+ ION, SAMPLING

OF KOH, % % AS KOH % AS KOH % AS KOH TECHNIQUE

CELL

NO.

BP2

BP3

BP4

BP6

BP7

BP8

BP10

26

26

31

36

21

26

23.5

25.5

23.6

28.0

33.2

17.3

24.03

21.6

5.27

2.7

3.3

2.5

4.0

5.07

5.1

30.77

26.3

31.3

35.7

21.3

29.1

26.7

SOXHLET*

BOTTOM OF CELL

BOTFOM OF CELL

BOTTOM OF CELL

BOT[OM OF CELL

SOXHLET*

SOXHLET"

* SOXHLET EXTRACTION OF #3 & 4 UNIT CELLS.

6

5

u_ 4

9 3

0 i I I I I I I I I I I ' , • i i i | I I • , , • • I

1000 10000 100000 1000000

NUMBER OF CYCLES

FIGURE 10. A PLOT OF CARBONATE BUILD-UP VS. NUMBER OF CYCLES.

TABLE 6, RESULTS OF CHEMICAL ANALYSES OF NEW AND CYCLED NICKEL

ELECTRODES FROM BP CELLS. NUMERICAL VALUES IN THIS TABLE

INDICATE AMOUNTS IN MMOLES/G OF ACTIVE MASS

NI° (SINTER)

TITRATION (A)

AA

M (ll&lll); TITRATION(B)

CO (ll&lll); AA

M(III)

TOTAL NI & CO

TITRATION

(A) + (B)

(A) l [(A),(B)], %

NEW BP1 BP2

27-08

8.29 4.92 3.70

8.20 4.59 3.27

4.90 7.24 8.10

0.67 0.55 0.57

1.50 0.90 1.78

13.40 12.06 11.94

13.1_ 12.16 11.80

62.85 40.46 31.36

BP4 BP6

7.79 8.41

6.4g 7.47

5.14 4.77

0.61 0.57

0.58 0,80

12,20 12,56

12.93 13.18

6O.25 63.81

BP8 BP10

4.99

4.59

7.32

0.57

2.60

12.23

12,31

40.54

4.59

4.2O

7,19

0.59

1.40

11.87

11.78

38.96

253
Chemical Analyses results of new and cycled nickel electrodes from test cells are shown in Table 6.

Active material content (amount of total ionic Ni and Co in nickel electrode) is plotted against numberof

cycles in Figure 11. This content increased with cycling indicating that nickel plaque may be corroded

gradually with cycling. Experimental data fit roughly to a linearcorrosion rate of 0.1 A/cycle (Curve A of

Fig. 11).

In order to see if Pt was carried over the nickel electrode from adjacent H2electrode in BP cells in

which soft shorts were developed. Elemental analyses (spectrographic) was carried out on outer edge

of nickel electrodes from a few cycled cells. Analysis results of a nickel electrode from BP9 showed a

considerable amount of Pt (0.25%) while Pt was not detected fro m a new electrode and detected only in

trace level (less than 0.01%)in electrodes from BP2 and BP5. Considering both BP2 and BP9 had soft

shorts, this non-unif0rm distribution of Pt might be dueto aviolent 02 recombination ("popping") at peripheral edge as a probable mechanism of Pt migration.

Scanning Electron Microscope (SEM)- SEM pictures of a cross-sectional view of cell stack compo- :

nents were taken as a part ofthe failure analyses. Sample pictures ofthe components from cycled cells

are shown in Figures 12, 13, and 14. BP6 (36%), which failed after relativelyshort cycling (1845 cycles),

showed no apparent sign of expansion or deformation of stack components. However, long cycled

cells, BP9 and BP2 (24,594 and 39,573 cycles, respectively) showed signs of heavy nickel electrode -

expansion, nickel active material migration, rupture of nickel sinter substrate, and deformation of other

.

components such as separator, H2electrode, and gas screen material. It is interesting to note that

despite the heavy expansion and sinter rupture in the long cycled electrodes, they suffered relatively ;

small capacity decrease indicating that these changes are not direct causes of the electrode failure. -"

8

;7

r_

F5

g

I I I

0 10O0O 2OOOO 3OOOO

NUMBER OF CYCLES

40000

FIGURE 11. A PLOTOF ACTIVE MATERiAL(30NTENT (TOTAL IONIC NICKEL AND

COBALT) OF NICKEL ELECTRODE VS NUMBER OF CYCLES.

±

254
DISCUSSIONS

Cycle lifeof a nickelelectrodeincreasedgreatly asthe KOHconcentrationdecreased inthe electrolyte, although initial capacity of a nickelelectrodedecreases slightly as the concentration decreases

reportedearlier_b.The cyclelifeincreasedmorethantwicewhentheconcentrationwasdecreasedfrom

36 to 31%. Although resultswithvery low KOHconcentration (21to 23,5%) was complicated due to

unusablelow voltage secondary plateauformationTM, the cycle life improved roughly nine times when

the concentration was reduced to 26% from conventional value of 31%.

The failure mechanisms of a Ni/H 2cell with a low KOH concentration (26% or lower) and one with a

high KOH concentration (31% or higher) appears to be different from each other. High concentration

cells failed on cycling by a gradual fading of nickel electrode capacity after relatively small number of

cycles (1845 to 4230), while low concentration cells failed on cycling after a large number of cycles

(mostly 25,000 to 40,000) by a "soft" short formation without an excessive capacity decrease of nickel

electrodes. As a result of the tremendous difference in the number of cycles, heavy electrode

expansion and active material migration and extrusion were observed in the long life low KOH

concentration cells, but relatively small physical changes are present in the failed nickel electrodes in

the high KOH concentration cells. However, we have observed much less "black powder" formation in

BP7 (21% KOH) after 9402 cycles than in BP4 (31% KOH), BP5 (31% KOH), and BP6 (36% KOH)

which were tested for less than half as many cycles as BP7, indicating that some physical changes

occur at much reduced rate in low KOH concentrations.

The observation that no significant internal resistance change of cells despite heavy electrode

expansion causing sinter rupture and active material migration and extrusion appears to indicate that

the conductivity of nickel sinter may not be the controlling factor in electrode failure. We speculate that

the capacity fading mechanism might involve crystallographic and micro morphological changes of

active material. All long life cells failed eventually on cycling due to a "soft" short formation without an

excessive capacity decrease of nickel electrodes as shown in Table 4. The positive identification of Pt

on a nickel electrode (spectrographic analysis result) appear to indicate that the short might be due to

the Pt on nickel electrode. This Pt might act as an hydrogen electrode while it is electrically shorted to

the nickel electrode through a layer of high resistance nickel oxide/hydroxide material. An alternate

mechanism might involve a short between nickel and hydrogen electrodes via extruded nickel active

material in the gas screen area.

255
ORIGINAL PAGE

BLACK AND V_HITE PHOTOGRAP.H

(a) (b)

FIGURE 12. SEM CROSS-SECTIONAL VIEW OF (A) #3 AND #4 CELLS AND (B) EXPANDED

VIEW OF LOCAL AREA OF BP6,

(a) (b)

FIGURE 13. SEM CROSS-SECTIONAL VIEW OF (A) #3 AND #4 CELLS AND (B) EXPANDED

VIEW OF LOCAL AREA OF BP9.

(a) (b)

FIGURE 14. SEM CROSS-SECTIONAL VIEW OF (A) #3 AND #4 CELLS AND (B) EXPANDED

VIEW OF LOCAL AREA OF BP2.

256 ORIGINAL PAGE iS

OF POOR QUALITY

=

I

e

i

=

=

z
ACKNOWLEDGMENT

We would like to acknowledge NASA-Lewis Research Center for its support for this study (Contract

NAS 3-22238; Project Manager: John Smithrick).

REFERENCES

. (a) H. S. Lim and S. A. Verzwyvelt, "KOH Concentration effect on the cycle life of nickel-hydrogen

cells." Proc. 20th IECEC, August 1985, p. 1.165. (b) H. S. Um and S. A. Verzwyvelt, "Cycle life of

nickel-hydrogen cells. II. Accelerated cycle life test, "Proc. 21 st IECEC, August 1986, p. 1601. (c)

H. S. Lira and S. A. Verzwyvelt, "KOH concentration effect on the cycle life of nickel-hydrogen

cells. II1.Cycle life test, "J. Power Sources, 22, 213 (1988).

. (a) H.S. Lim, "Long life nickel electrodes for nickel-hydrogen cells", Final Report of Phase I, NAS

3-22238, NASA Cr-174815. December, 1984; (b) H.S. Lim and S. A Verzwyvelt, "Long life nickel

electrodes for a nickel-hydrogen cell: I. Initial Performance, "Proc. 18th IECEC, August 1983, p.

1543; (c) H.S. Lim and S. A Verzwyvelt, "Long life nickel electrodes for a nickel-hydrogen cell:

Cycle life test," Proc. 31st Power Sources symposium, June 1984, Cherry Hill, NJ. p. 157; (d) H.S.

Lim and S. A Verzwyvelt "Long life nickel electrodes for a nickel-hydrogen cell: Results of an

accelerated test and failure analyses. "Proc. 19th IECEC, August 1984, p. 312

. H. S. Lim and S. A. Verzwyvelt, ' Nickel electrode expansion and the effect of LiOH additive." Proc.

20th IECEC, August 1985, p. 1.104.

. M. P. Bernhardt and D. W. Mauer, Proc. 29th Power Sources Conf. 1980, Electrochemical Society, p. 219.

. P.P. McDermott, "Analysis of Nickel Electrode Behaviorin an Accelerated Test." Proceedings of

Symposium on the Nickel Electrode, Proc. Vol. 82-4, R. G. Gunther and S. Gross Eds. (The

Electrochemical Society, Pennington, NJ, 1982), p. 224.

V. A. Tracy, "The properties and some applictions of carbonyl-nickel powdwers," Powder

Metallurgy, 9, 54 (1966).

257
2

2 2_ 2 L L
N90-20476

ADVANCES IN LIGHTWEIGHT NICKEL ELECTRODE TECHNOLOGY

Dwaine Coates, Gary Paul and Paul Daugherty

Eagle-Picher Industries, Inc.

Electronics Division

Joplin, Missouri 64802

Studies are currently underway to further the development of lightweight

nickel electrode technology. Work is focused primarily on the space nickelhydrogen system and nickel-iron system but is also applicable to the nickelcadmium and nickel-zinc systems. The goal is to reduce electrode weight while

maintaining or improving performance thereby increasing electrode energy

density. Two basic electrode structures are being investigated. The first is

the traditional nickel sponge produced from sintered nickel-carbonyl powder.

The second is a new material for this application which consists of a non-woven

mat of nickel fiber. Electrodes are being manufactured, tested, and evaluated

at the electrode and cell level.

INTRODUCTION

Eagle-Picher has been involved in nickel electrode development for a

number of years in association with many different battery systems. Our

current work with the nickel-hydrogen system dates back to 1969 and has evolved

from previous work with aerospace nickel-cadmium and zinc-air battery cells.

The goal throughout has been to increase the energy density of the nickel

positive electrode (and hence the battery) while maintaining a satisfactory

level of performance. For space applications, performance includes sufficient

useful life as well as cell electrical characteristics.

The goal of increasing cell energy density provided the original impetus

for the development of the nickel-hydrogen system. Replacing the heavy cadmium

negative electrode of the nickel-cadmium cell with a lightweight catalytic gas

electrode provided a substantial savings in weight. Nickel-hydrogen cell

designs have been continuously refined over the years to remove even more system

weight. However, cell pressure vessels and structural components can be reduced

in weight only to a finite degree due to safety considerations and launch dynamics requirements. Common pressure vessel battery designs hold some promise

for further weight reductions but are not currently being used in production

programs.

The best opportunity for further weight savings without serious deviation

from existing flight cell designs would appear to lie in the electrochemically

PRECEDING PAGE BLANK NOT FILMED

259
active cell components. The nickel electrodes account for about forty percent of

the total finished weight of a current design flight production nickel-hydrogen

cell. Therefore even a small increase in the energy density of the electrode

would amount to a substantial benefit at the cell and battery level.

The nickel electrode consists of two basic components. These are the

active material and the support matrix. Active material utilization has been

optimized in past efforts and probably can't be increased muchpast current

levels. The most likely opportunity for weight reduction is in the structural

matrix of the electrode which houses the active material. For conventional

sintered nickel-carbonyl powder electrodes this would require producing a higher

porosity matrix, a lighter weight current collector/substrate, or a thicker

electrode. Use of a nonwovenmat of nickel fibers as an alternative lightweight, high porosity matrix is also being explored.

SINTEREDNICKELCARBONYLPOWDERELECTRODES

The production of battery electrodes from sintered nickel carbonyl powder

is a well established technology and has been extensively investigated and

reviewed. However, work is still being done to further refine the process for

high reliability, long life, and weight-critical applications. Twobasic

processes are currently in use to produce electrodes for the nickel-hydrogen

system. One is a dry powdermethod in which each individual plaque is hand-laid

in a mold one at a time. The second is a wet slurry process in which the plaque

is produced in a continuous roll. Both methods use a similar nickel wire mesh

substrate and sintering process and yield a similar product.

=

Once the porous nickel matrix is obtained it must be loaded with active

material to produce a working electrode. The chemical impregnation method used

for many years to produce nickel-cadmium electrodes has been entirely replaced

by an electrochefnical process in the nickel-hydrogen system. Electrodes are

produced in either an aqueous solution or an aqueous/alcohol solvent system.

Both processes have the advantage of loading the active material i,ione

impregnation step rather than many repeated cycles as in the chemical process.

Other advantages of electrochemical impregnation include improved cycle life,

utilization, and dimensional stability.

Electrodes produced by the wet slurry method and the dry powder method

exhibit relative advantages to each other. Both processes are characterized by

extensive quality control provisions (e.g. strict, tight tolerance requirements

from raw materials, operational process steps and finished component testing) in

recognition of their criticality to the finished product. The slurry process is

a continuous-feed, more efficient, and less operator-dependent process. The

current-collector substrate is automatically centered in the electrode structure

by the nature of the process. This provides electrical and structural benefits.

The material has high mechanical strength in terms of both tensile and bend properties. Slurry type plaque hold Up to the rigors of electrode processing and

handling resulting in a low attrition rate during the multi-step production process. The slurry type material better accepts attachment of electrical conductors by resistance welding without forming cracks in the heat affected zone and

surrounding sinter. Part of the slurry production process involves a wet oxidation step which covers the surface of the plaque with a thin passive oxide

260
layer. This serves to provide a more durable product with low susceptibility to

mechanical damage or chemical corrosion. Most importantly, slurry type electrodes exhibit good dimensional stability and should provide a longer cycle life

at higher depths-of-discharge.

Electrodes produced by the dry powder method possess inherent battery

weight and energy density advantages. Dry powder process electrodes have a

higher porosity structural matrix. This results in a higher active material to

inactive substrate weight ratio and greater energy density. Dry powder electrodes also are manufactured with a greater electrode thickness. Increasing the

thickness of the electrode has a net effect of increasing cell energy density by

requiring fewer electrodes and associated cell components to achieve the same

ampere-hour capacity.

Historically, applications of the space nickel-hydrogen system have been

with slurry type electrodes impregnated by the aqueous electrochemical method.

This extensive data base has been previously reviewed (Reference 2) and includes

five major satellite programs (18 launches total) which are currently operating

in orbit with nickel hydrogen batteries on board. The ideal electrode, therefore would be one which combines the mechanical strength and dimensional stability of slurry process plates with the higher porosity and greater electrode

thickness associated with dry powder process plates.

Accordingly, manufacturing parameters were adjusted to produce several lots

of 35 mil thick slurry process plaque with porosities ranging up to 86%. These

plaque retained 60 to 80 percent of the nominal mechanical strength of normal

30% porosity, 30 mil thickness slurry plaque. Each lot was impregnated and

fonned by normal manufacturing procedures, and two plates from each lot were

subjected to a standard flooded capacity measurement and a 200 cycle stress

test. _l plates tested exhibited satisfactory utilization numbers in the 120

to 130% range. However, a growth versus porosity plot of the 200 cycle, I0C

stress test data (Figure 1) shows that there is currently a breakover point at

around 83% porosity beyond which the amount of electrode growth becomes

excessive. Thickness growth is undesirable as it has been previously identified

as a major contributor in long term cycle failure of nickel-hydrogen cells

(Reference I).

Two lots with porosities of 82.7% and 84% were selected for further

cycling. These plates underwent a total of I000 stress test cycles with a

visual inspection for physical integrity and growth measurements conducted at

the completion of 200, 600, and I000 cycles. Flooded capacity tests were run at

the conclusion of each 200 cycle increment. None of the plates blistered or

suffered a loss of capacity during the I000 cycles of testing. Thickness

increase peaked at 600 cycles for both lots. The 82.7% porosity lot exhibited

less than 6% thickness increase after the completion of I000 cycles.

Figure 2 shows the relationship between electrode growth and active

material loading on a grams per cubic centimeter of void volume basis. This 200

cycle 10C stress test data suggests that the loading of high porosity slurry

plates can be increased beyond current flight acceptable levels and thereby

increase the energy density.

The goal of future development work is to push the breakover point

demonstrated in Figure I beyond 84% porosity through further refinement of

261
slurry process manufacturing parameters. Impregnation and formation of these

high porosity slurry plaque will take place in Eagle-Picher's new nickelelectrode production facility. This facility has been expressly designed and

dedicated to the space nickel-hydrogen system. As noted above, in recognition

of the criticality of the component, the facility was _designed from the group up

for full automated computer process control. The generated process data forms

the basis for a quaiity control statistical analysis with "real-time" feed-back

provisions offering continuous, detailed process control.

Improved tank flow characteristics, greater solution volume to plaque area

ratios, automated cycling and pH control, an oil-cooled balancing resistor network for equal current distribution, and computerized monitoring and recording

of production parameters will allow tighter control and reproducibility of the

manufacturing process. Slurry process electrodes with nominal porosities of 82%

have already been produced in Eagle-Picher's existing facility, evaluated at the

electrode level, and are currently being assembled into cells. It is anticipated a nominal mass saving of approximately 10% will be achieved at the

finished cell level.

NON-WOVEN NICKEL FIBER ELECTRODES

Work has been done over the last_several years on a non-woven nickel fiber

material as a possible replacement for conventional sintered nickel powder type

plaque. The material is manufactured by the National Standard Company under the

trade name Fibrex (TM) and has received renewed interest recently. It has the

potential for substantial improvements in energy density and material cost

reduction. Some of the previous work with fiber mat had been unsuccessful

because of problems in loading the activematerial a_d ther_ui_ng dimensional

instability of the finished nickel electrode. The major problems appeared to be

due to the open structure and large pore size of the mat and its lack of mechanical strength.

More recent work has been conducted mostly in relation to the nickel-iron

battery system under the Dual Shaft Electric Propulsion (DSEP) Program funded by

the United States Department of Energy (DOE). This investigation has been undertaken due to the development of a nickel fiber mat which has been cross-linked

with battery grade nickel powder. This material has the potential for providing

the necessary mechanical strength and dimensional stability required for long

cycle life battery electrodes. Current work is focused on adjusting the fiber

to powder ratio and the sintering conditions to achieve an optimum electrode for

a given cell design. Cross_sectionai strengths have been increased by a factor

of ten over the earlier materials. The fiber mat promises good performance and

design versatility through close control of electrode parameters such as weight,

strength, surface area, pore size, and thickness. A 12% savings in weight over

the sintered powder electrode is realized solely due to the elimination of the

grid wire support/current collector structure not required by the mat. Further

weight savings may be achieved from the higher porosities and active material

loading obtainable from the fiber mat.

The problems associated with loading the active material into the fiber mat

have been overcome in part by applying a periodic reverse impregnation process.

The method is essentially the same as the normal aqueous electrochemical process

262
except that the polarity of the electrodes is periodically reversed during

impregnation. Periodic current reversal is commonlyused in the metal plating

industry to promote a more uniform deposited film on the work piece. It appears

to have a similar load leveling effect in impregnating nickel electrodes with

active material. Work is underwayto continue investigation of plating/

deplating current densities and cycle times to achieve optimum results. This

method has an additional benefit of yielding a 20%higher ampere-hour efficiency

in loading the plate over the normal constant current process.

Nickel-iron cells have been constructed and tested using electrodes produced from a 50%fiber/50% powdernickel fiber mat. The porosity of the plates

was 75%, the thickness 0.040 inches and the nickel fiber diameter 20 microns.

The plates were impregnated through two cycles due to inadequate active material

weight gain after the first cycle. After formation in aqueous potassium

hydroxide the plates had an average active material loading of 1.54 grams per

cubic centimeter of void volume (g/ccvv). A plate was assembled into a test

cell and 50 cycles were performed at 100%DOD(Figure 3). On-going research is

now exploring sinter porosities up to 85 to 90%which will offer specific

energy advantages.

Additional testing was developed for use in the DSEPProgram to determine

the best electrode design for optimum power capabilities. The test is composed

of a base discharge at a C/2 rate with 2._5C rate pulses at 95%, 75%, 50%, and

20%of the working capacity. Testing was performed on fiber electrodes (Figure

4) with normal sintered electrodes as a control. It was concluded that the

sinter electrode gives greater instant power within the first three pulses.

However, the sinter electrode appears to increase in impedanceover the

discharge cycle more than the fiber electrode. The fiber electrode also yields

a greater utilization of active material per gram than the sinter electrode. The

fiber electrode could be developed to supply more power over the discharge curve

by improving its dimensional stability. This would increase the contact area of

the fibers thus lowering the internal resistance of the plate.

Further investigation into the utilization and electrochemical efficiency

of the fiber electrode was madeby Dr. Alvin J. Salkind and Associates. This

was accomplished by a subcontract from the base technology contract with DOEfor

the development of the nickel-iron battery.

Testing was performed comparing the rate capability of 0.063 in. thick

sintered and fiber nickel plates at various states of charge (SOC) but the

fiber-based plate supported higher current densities at the lower SOC of 25%.

The impregnation loading level of the fiber nickel plate was much lower than the

standard sintered plate, with a consequent reduction of capacity.

The marked differences in scope and ability to sustain heavy loads at a low

state of charge may be only a consequence of the relatively low active material

content of the fiber plate. However, it seems likely that the fiber material

can be loaded higher and still perform as well as a sintered plate of comparable

active material density.

263
CONCLUSION

Considerable progress has been made toward improving the nickel electrode

for high reliability, long cycle life space applications. The goal has been to

increase energy density while maintaining currently acceptable flight program

active material loading levels. In addition, the goal has been to significantly

advance the level of quality control provision implementation and the level of

quality control process data utilization for enhanced, "real-time", total process control. Electrodes produced for early nickel-hydrogen flight cell

programs had an average energy density of about 0.135 ampere-hours per gram.

Current flight production programs have achieved energy densities at the

electrode level more than 20% greater than that value at the same active

material loading on a grams per cubic Centimeter of void volume basis. This

translates into a substantial weight savings at the finished flight iev#l.

Further increases should be attainable in the near future within the scope

of current production technology. The main thrust of Eagle-Picher,s future

research work will be to develop a mechanically strong, thick, high porosity

nickel electrode. Current data suggests that this electrode will more likely be

produced using the wet slurry manufacturing process; but, work will continue

with both the slurry and dry powder processes.

Nickel fiber mat appears to be a promising replacement for the traditional

nickel sinter substrate. Preliminary testing has shown it to be a viable battery electrode with distinct weight advantages. Work will be continued towards

the further development of optimized fiber electrodes for nickel battery

systems.

In add_t_o_n_ to superior physical performance and enhanced component

quality, it is anticipated both of the above electrode substrate materials will

eventually result in overall system cost reductions.

ACKNOWLEDGEMENTS

The authors wish to acknowledge the Department of Energy Funding of the Dual

Shaft Electric Propulsion Program and the National Standard Company for providing fiber mat material in support of the non-woven nickei fiber electrode

work. Much of the early basic testing and evaluation work on nickel sinter was

performed-by the COmmunication Satellite Corporation Laboratories funded by the

International Telecommunications Satellite Organization.

REFERENCES

. Coates, D.K., Barnett, R.M.: Nickel-Hydrogen Cell Life Testing, 23rd

Intersociety Energy Conversion Engineering Conference, Volume 2, 1988,

pp. 483-488.

Miller, L.: The Ni-H 2 Battery System: A Space Flight Application

Summary, 23rd Intersociety Energy Conversion Engineering Conference,

Volume 2, 1988, pp. 489-492.

264
14.0

13.0

12.0

uJ II.0

m I0.0

g.o

U

z 8.0

I.,,,,,4

7.0

8.0 5.0

1..1

_ 4.0 T

_- 3.0

2.0

1.0

0.0

79

ELECTRODE

High

¥S POROS 1]T1{

Slur rLj

0

O'

0 o

/

0

l I I I I I I

80 81 82 83 84 85 88

POROSITY

FIGURE I

87

14.0

13.0

12.0

11.o

IO.rl

9.0

8.0

7.0

_ 8.0

g s.o

4.0

_c 3.0

2.0

1.0

0.0

1.50

ELECTRODE

High

_RO_TH

P0r0s itg

VS LOAD 11INS

Slurrg

o

o

o

o

I I I l I 1 L.

i.55 1.60 i.85 1.70 1.75 1.80 1.85

LDFIDING (6LCCVV)

FIGURE 2

!

1.90

I

1.95 _.00

265
FIBER ELECTRODE PULSE RATE TEST

1.50

1.35

LU

LO

(:

1.06

0.90

0.?5

0

\

\

I ! I | I I I |

I0 2O 3O 40 50 8O 70 80 90

TIME (MIN.)

FIGURE 3

!

IO0 IIO

25

2O

_o

15

0

D

FIIBER ELECTRODE CYCLE TEST

IL,

IDO

I !

2OO

IO0 Z 000 CYCLES

FIGLI_E 4 _:

266

500

z
W

(.3

II

ii!

I--

>

! .40

1.35

1.30

1.25

1.20

1.15

!.10

1.05

1.00

0.0

F IBER Iq! CI(EL ELECTROI]E

YI)LTASE ¥S. IDgSCHARGn::RATIE 6 SO(;

[] .... __ ...... _._

.....

I I I ., [ ,, I I -L-----,_. I I J

0.5 [.0 1.5 2.0 2.6 3.0 3.5 4.0 4.5 5.0

O ISCHF_GE RRTE

FIGURE 5

1.4

1.3

1.2

1.I

1.0

o.g

0.8

O.O

S _qTERED _ _CKEL ELECTRODE

YOLT_G_ YS. _OSCH_R6E R_IP_ 8 SQC

_ I | I I z _ _1,_ I .,7

0.5 ! .O 1.5 2.0 2.5 3.0 3.5 4.O 4.5 5.0

DI SO-IFIRGE F_TE

FIGURE 0

267
m

I
N90: 0477

MULTI-MISSION Ni-H 2 BATTERY CELLS FOR THE 1990's

Lee Miller, Jack Brill and Gary Dodson

Eagle-Picher Industries, Inc.

Joplin, Missouri 64802

A sufficient production, test and operational database is now available to

permit design technology optimization for the next decade. The evolved battery cell

design features standardized technology intended to support multiple type missions

(e.g. both GEO and LEO). Design analysis and validation test cells demonstrate

that improved performance plus attractive specific-energy characteristics will be

achieved.

INTRODUCTION

In the year 1988 Eagle-Picher Industries (EPI) surpassed the production point

of I0,000 space type Ni-H 2 battery cells. Approximately 25,000,000 battery cell

hours of space flight operation have now been accumulated. Multiple test cell

groups continue to undergo real-time and accelerated life testing. Real-time CEO

testing has now exceeded 13 years and LEO testing over six (6) years. Accelerated

LEO testing has exceeded 40,000 cycles.

Several design technologies evolved in the accumulation of this extensive

data-base. Performance comparisons and post-test destructive physical analysis

(D_A) have proven useful in the evaluation of the advantages and disadvantages of

the various design technologies. Sufficient data is therefore available to propose

a more standard, optimized battery cell design capable of supporting multiple types

of missions for the next decade.

BATTERY CELL DESIGN

The following design summary proposes a baseline cell design. If optional

technology is available which has been tested and proven, then it is noted in parenthesis.

Pressure Vessel (pv)

A single girth weld design is proposed with opposing axial compression terminal

seals as shown in Figure I. (For volume critical or other special applications,

the terminals may be located at 45 ° off-set positions on either the same PV end or

at opposing ends.)

In accordance with MIL-STD-1522A, EPI has established and qualified an "inhouse" fracture critical vessel inspection capability which permits customizing

the cell design maximum-expectcd-operating-pressure (MEOP) and vessel wall thickness for the specific mission application.

PRECEDING PAGE BLANK NOT F;LMED

269
ORIGINAL PAGE IS

OF POOR QUALITY

Electron Beam Weldin$

All PV joining will be accomplished by the electron beam (EB) process. EPI

has procurred a FERRANTI Sciaky 60 KV, 250 mA EB welder for "in-house" capability

and control. A large vacuum chamber was selected to permit equipment modifications

for high precision, multiple cell self-indexing, fully automated welding.

Electrode Stack

Z

The electrode stack will feature "pineapple slice" geometry for optimum

thermal, mechanical and electrical characteristics (refer tO Figure I). The positive electrodes will be configured in a "back-to-back" arrangement. Depending upon

capacity the design will feature a single or dual stack arrangement. =or capacities

of 50 Ampere-Hours (A-H) or less, a single stack is generally used.

PV Wall Coating

The PV will feature a porous zirconium oxide wall coating to serve as an

electrolyte return (wick) and reservoir. (A combination electrolyte transport and

catalyzed wall coating design for enhanced gas management is also offered.)

Positive Electrode

A high mechanical strength (1500 psi minimum), slurry sinter positive electrode

is utilized for long term dimensional stability. An electrode thickness of 0.03

inches (0_076 cm) is incorporated to maximize the quantity of electrodes and thus

minimize operational current densities. A moderate active material loading level

of 1.65 grams/cm 3 of void volume is specified.

Negative Electrode

A standard teflonated catalyzed electrode with a platinum loading of 8 milligrams/cm 2 is featured. (Platinum loading reductions of up to an order of magnitude

are also available).

Separator

In consideration of long term availability, electrolyte reserve provisions and

low impedance (particularly at lower temperatures), a two (2) layer zirconium oxide

(Zircar) material design is featured. (A combination asbestos/Zircar design is

also available. In addition, for mass critical missions a single laver asbestos

or Zircar design is offered.)

Electrolyte

A 31% potassium h,!droxide solution will be incorporated in the standard cell.

(rot missions with a large number of Cycles or high depths of discharge, a 26_

solution is offered for nositive electrode _tress reduction.)

Gas ]_ana_ement ; : _

The more open structural characteristics of the Zircar senarator material

poses additional problems with respect to oxveen _as management (reference =i_ure

2). The proposed cell desien will incorporate provisions which are intended to

270
ORIGINAL PAGE IS

OF POOR QUALITY

redirect the oxygen gas flow away from the positive electrode-to-senarator interface

during overcharge.

Stack Growth Accommodations

To accommodate potential electrode stack dimensional _rowth and to maintain

appropriate stack load under dynamic stress, a sprin_ type device is employed. The

device incorporated in the proposed cell design offers significant advantaEes over

the Belleville washer design employed in several cell types. The spring constant

for the proposed device can be reproducibly controlled and maintained in the manufacturing process. In addition, load uniformity is maintained over much greater

length of travel at approximately one fourth the mass of equivalent Belleville

washers.

80 A-H CELL DESIGNS

Battery cells rated at 80 A-H are being produced and validation cells have been

placed on test. This program may serve as an example for projecting mass and volume

characteristics for the proposed cell design. Figures 3 through 6 represent summaries of computer projections for the most influential technology option, separator

type and design.

The design analysis of the proposed Ni-H 2 battery cell was performed on a TRS-80,

Model i00 portable computer. The specific application software was coded in an extended version of BASIC. The program constructs a detailed model of the cell design

via established electrochemical, physical and material performance formulas.

The 80 A-H cells for a program have now progressed sufficiently to permit model

accuracy verification.

Attribute Model Accuracy

Mass 0.2,%

Capacity 0.2_

Length 0.1%

CELL TEST DATA

Typical battery cell conditionin_ and Accemtance Test data are presented in

Tables 1 and 2, respectively. To permit a level of performance comparison, data was

selected for a single layer asbestos and a sln$1e layer Zircar separator cell design.

These same two (2) test cell groups are mart of a larger test cell Eronp now

undergoing real-time GEO life cycle testin_ in accordance with the test profile

presented in _i_ure 7. Typical end-of-discharge voltage (EODV) performance is

compared _raphically in _i_ure 8.

CNNCLUSION

A multi-mission Ni-H 2 battery cell has been oroduced and is successfully under-

_oin_ validation testing. The cell design has assimilated the o_timal, space flight

proven technology which has evolved from a 15 year. I0,_00 mroduction unit,

25,000,000 flight cell hour database. Its standardized features should now sup_or_

missions through the 1990's with minimal need for de_i_n requalification and dedicated cycle life testing.

271
TABLE 1.

_ycle

1

2

3

Notes:

TABLE 2.

" L__ZesL

25°C

_5°C

I0°C

Notes!

Performance Comparison of 80 Ampere-Hour Cells

During Sealed Conditioning Cycles*

EOC V

#1 ! #2

1.535 1.534

1.5271 1.525

1.537! 1.533

t

1 HrDV

#1 #9

1.22O 1.233

1.218 1.233

1.215 1.234

#1 _ #2

88 86

85 85

84 85

* Charged C/10 for 16 hours, discharged C/2, I0°C.

1. Design #1 = Asbestos Separator, 8 cells on test.

2. Design #2 ---Zircar (1) Separator, 3 cells on test.

3. All data values are group averages.

Performance Comparison of 80 Ampere-Hour Cells

During Acceptance Test Cycles*

' F_QC__L

i #! #2_

i

1.491 1.486 i

_ 1.595 1.585

1.530 1.515

i

1 HrDV ' _Cap. (A-H)

1.204 J 1.211 _ 73 1 71

1.140;1.188182!84

1.185 j 1.205 , 87 _ 86

, i , (77)**i (84)** I i i

* Charged C/i0 for 16 hours, discharged C/1.6.

** Discharge to 1.1 volts.

1. Design #1 = AsSe_6s Separator, 8Celi_on test.

2 Design _2 = Zircar (t) Separator, 3 cells on test.

3. All data values are group averages.

4. Cells were subjected to 30 "burn-in" cycles between

Conditioning and Acceptance Testing. Cycle = 12 hour;

75% DOD in t.2 hours, 1.25 C/D return in 10.8 hours,

10oc.

=

272 __
_ f

,.llllilllll_ _

SECTION B-B

Figure 1. EPI "Mantech" Design
CONVENTIONAL

O2

FLOW

POSITIVE

o_

tll I11 II111111[t{ { I1{

NEGATIVE 1

ASBESTOS

, _ =-:

CONVENTIONAL

O2

FLOW

EPI REDIRECTED

1⁄2

FLOW

o_

POSITIVE 1

_, o_

ili ,_ililj,,lilt_i _ j_t

NEGATIVE t

--- POSiTiVE

i i i I i I i i i i i i

NEGATIVE

2 LAYER

ZlRCAR

SINGLE LAYER

ZIRCAR

Figure 2. NiH 2 Separator Design Technology

274
Figure 3.

Value_____: Batte__L_JLDe__g n Parameters (_n_utL"

1. Rated Cell Capacity (A-Hr)

2. Cell Capacity Margin (%)

3. Cell Residual Capacity (%)

o

a_/Cell Desi F_s O_Q__p_ut_;_

1. Nominal Cell Mass (Gms)

2. Nominal Cell Capacity (A-Hr)

3, Cell Specific Energy 0N-Hr/Kg)

4. Pressure Vessel Length (in)

5. Total Cell Length (in)

6. Total Cell Length (cm)

7. Cell Diameter (in)

8. Cell Diameter (cm)

!

80.000

10.000

25.000

1991.280

88.476

54,207

10,123

13.373

33.968

3.506

8,905

_J

Baseline 80AH Rated (2 Layer Zircar Separator) Cell Design

Figure 4.

Batte!3/_cell Design Parameter s_(_n_pu_ Value:

1. Rated Cell Capacity (A-Hr)

2. Cell Capacity Margin (%)

3. Cell Residual Capacity (%)

Batte_ Cell Desi_ature=s,_LO_u_tp_,u_tJ;

1. Nominal Cell Mass (Gms)

2. Nominal Cell Capacity (A-Hr)

3. Cell Specific Energy (W-Hr/Kg)

4. Pressure Vessel Length (in)

5, Total Cell Length (in)

6, Total Cell Length (cm)

7. Cell Diameter (in)

8. Cell Diameter (cm)

o

80.000

10.000

25.000

Y aLu_e:

1947.790

88.476

55.417

10.021

13.271

33.707

3,506

8.905

Option =1 80AH Rated (Asbestos / Zircar Separator) Cell l-)r.si_.=

275
I B.attery Cell Design Parameters (Input): Value:

1. Rated Cell Capacity (A-Hr) 80.000

2. Cell Capacity Margin (%) 10.000 I

3. Cell Residual Capacity (%) 25.000

Battery. Cell Design Features (Output); Value:

1858.410

89.271

1. Nominal Cell Mass (Gms)

2. Nominal Cell Capacity (A-Hr)

3. Cell Specific Energy (W-Hr/Kg) 58,6_04

4. Pressure Vessel Length (In) 9.643

5. Total Cell Length (in) 12.893

6. Total Cell Length (cm) 32.749

7. Cell Diameter (in) 3.506

8. Cell Diameter (cm) 8.905

!

I )

J

I

Figure 5. Option #2 80AH Rated (1 Layer Asbestos Separator) Cell Design

Battery Cell Design Parameters (Input):

1. Rated Cell Capacity (A-Hr)

2. Cell Capacity Margin (%)

3. Cell Residual Capacity (%)

Battery. Cell Design Features (Output);

1. Nominal Cell Mass (Gms)

2. Nominal Cell Capacity (A-Hr)

3. Cell Specific Energy (W-Hr/Kg)

4. Pressure Vessel Length (in)

5. Total Cell Length (in)

6. Total Cell Length (cm)

7. Cell Diameter (in)

8. Cell Diameter (cm)

!

Value: !

80.000

10.000

25.000

Value:

1803.090

88.476

59.865

9.505

12.755

32.397

3.506 l

8.905

.2

Figure 6. _Opt-ion #3 80AH Rated (i Layer Zircar separator)Cell Design

=

=

|

276 h

z
FIGURE 7. Cycle Test, Cell Temperature Vesus Time

T

E

M

P

25°C

.5oc

..................... ,; ......... : i:\ -A DODK i-

/!\

I I

I I

I I

I I

INITIAL 20 HR CHARGEOR 17 HR ',TRICKLE', 1.2 HR

CHARGE PLUS TRICKLECHARGE ,CHARGE, DI_HG'

4 .-,-...... _',<---_¢_i

I

75 %

1.15

24 Hrs

17 HOUR CHARGE PLUS TRICKLE

CHARGE

k

T

TIME

E

O

D

V

o

I

t

R

g

e

FIGURE

1.20

1.19

t.18

1.17

1.16

1.15

1.14

1.13 .....

0

8. Real-time GEO Life Cycle Test End-of-Discharge Voltage

Design #2, Zircar

Design #1, Asbestos DOD = 75 %

K = 1.15

Cycle - 24 Hours

I i

5

,I i i __1_

10 t5 20

Cycle Number

25 30

277
T7
N90-20478

NICKEL-HYDROGEN CAPACITY LOSS ON STORAGE

Michelle A. Manzo

NASA Lewis Research Center

Cleveland, Ohio 44135

A controlled experiment evaluating the capacity loss experienced

by nickel electrodes stored under various conditions of

temperature, hydrogen pressure, and electrolyte concentration was

conducted using nickel electrodes from four different

manufacturers. It was found that capacity loss varied with

respect to hydrogen pressure, and storage temperature as well as

with respect to electrode manufacturing processes. Impedance

characteristics were monitored and found to be indicative of

electrode manufacturing processes and capacity loss. Cell

testing to evaluate state-of-charge effects on capacity loss were

inconclusive as no loss was sustained by the cells tested in this

experiment.

INTRODUCTION

Nickel-hydrogen batteries are rapidly becoming accepted for use

in low-earth-orbit and geosynchronous orbit applications. With

their increased use, it has become evident that the storage

procedures commonly used for nickel-cadmium cells are not

adequate for the nickel-hydrogen system. Frequently cells that

have been stored, under varied conditions, have exhibited a

significant loss of capacity. However, much of the data available

on capacity loss has been confusing and often conflicting. (ref.

1-6) As a result, members of the Electrochemical Technology

Branch at the NASA Lewis Research Center conducted a controll_d

determination of the capacity loss exhibited by nickel elec_rua=s

when exposed to different storage conditions.

A comprehensive test matrix was developed to evaluate the

capacity loss characteristics of nickel electrodes from four

different manufacturers. Two types of tests were run; individu_l

electrode tests, which involved flooded capacity and impedance

measurements before and after storage under varied conditions of

temperature, hydrogen pressure, and electrolyte concentration;

and cell tests which primarily evaluated the effects of state of

charge on storage. The electrode tests were run using electrod_

from Gates Aerospace Batteries (GAB), Whittaker-Yardney (W/Y),

Hughes (HAC), and Eagle Picher-Joplin (EP). The cell tests were

run only with Hughes electrodes. The cell tests evaluated

capacity loss on cells stored open circuit, shorted and trickle

charged at C/200 following a full charge.

PRECEDING PAGE BLANK NOT FILMED

279
The results indicate that capacity loss varies with the specific

electrode manufacturing process, storage temperature and storage

pressure. In general, electrodes stored at low temperatures or

low hydrogen pressures exhibited a smaller loss in capacity over

the twenty-eight day storage period than those stored at high

pressure or high temperature. The capacity loss appears to

correlate with the level of cobalt in the nickel electrode, with

a greater capacity loss observed in electrodes with higher cobalt

levels. Impedance measurements appear to correlate well with

capacity loss for a given type of e!ectrode but do not correlate

well with the capacity lPgs_pb_ve_w_e_o_es

fabricated by different manufacturers. There was a definite

correlation between the electrode potential measured immediately

following storage and the measured capacity loss. The C_l!_

storage tests indicated a temporary Ca pacity--ioss in_cellS - Stored

discharged in the open-circuit condition and shorted. The cell

stored on trickle charge did not exhibit a loss of capacity.

EXPERIMENTAL

NICKEL ELECTRODE STORAGE TEST PROCEDURES

Before Storaqe Flooded Capacity Measurements

Fifteen electrodes were obtained from each of four nickelhydrogen cell manufacturers for the capacity loss on storage test

matrix. The following manufacturers supplied electrodes for this

study: Eagle Picher-Joplin, Hughes, Gates Aerospace Batteries,

and Whittaker/Yardney. The characteristics of the electrodes are

summarized in Table i. The electrodes were divided, by

manufacturer, into two groups of thirty each for the flooded

capacity measurements. Three electrodes from each manufacturer

were subjected to the same set 0f_tes_:and storage conditions.

The initial capacity was estimated for electrodes from each

manufacturer. Where Capacity d_fferences existed between

electrodes in the groups £es£ed toget_e_r, _ test parameters (charge

and/or discharge rates) were adjusted to accommodate the sub

groups of electrodes. Since electrodes of varying capacity were

tested in the same series string, relative "C" ra£es may have

varied; however, in all cases the before storage and af_2

storage_measurgm@nt _ were made in the same way. (The C rate will

discharge the Domxnal electrode capacxty xn one hour.) Slnce £he

variable of interest was percent change in capacity, _£ was felt

that the results would be valid and comparable.

The flooded capacity measurements were made using a hickel

counter electrode and an amalgamated zinc reference electrode.

31% KOH was used for the majority of measurements. One group

from each manufacturer was tested in 26% KOH for a measure of

capacity retention vs electrolyte concentration. The initial

formation charges were run at C/7.5 or C/10 for 16 hours. The

28O
electrodes were discharged at the C/4 or C/5 rate. Formation

cycles were run until a uniform capacity was obtained on

consecutive cycles. This usually took six to seven cycles for

each group of electrodes. An average formation capacity was

obtained for each electrode.

Based on the formation capacity, C/2 charge and C/4 discharge

cycles, with a C:D of 1.12 to 1.15, were planned for the

electrode capacity measurements. However, as stated above, the

grouping of electrodes of unequal capacities lead to relative "C"

rates varying from C/4 to C/5.25 on discharge. A total of five

to six capacity determination cycles were run. An average before

storage capacity was determined for each electrode.

Before Storaqe Impedance Testing

Following the pre-storage capacity measurements the electrodes

were left in KOH in the discharged state until impedance

measurements could be made. A Princeton Applied Research 5208

Two-Phase Lock In Analyzer coupled with a Princeton Applied

Research Potentiostat-Galvanostat Model 273 was used to make the

impedance measurements. Impedance tests were run from .0007 Hz

to 5.65 Hz. All electrodes were tested in the discharged state,

following equilibration at 0.2 volts for i hour vs a Hg/HgO

reference electrode.

One electrode from each group of three was also measured in the

charged state. The electrode was charged at approximately the

C/3 rate for 1/2 hour, then equilibrated at 0.4 volts vs Hg/HgO

for 1 hour before measuring impedance. Following the impedance

measurements, the electrodes were returned to 0.2 volts by

discharging at rates similar to those used in the capacity

measurement cycles and equilibrated there for 1 hour before

storage.

Storage Procedures

As testing on groups of six electrodes that were to be subjected

to the same storage conditions, (three from each manufacturer in

the sub group) was completed, the electrodes were installed in a

vessel and put on storage stand for twenty-eight days. Storage

conditions and their identifying codes are described in Table 2.

Storage temperature, hydrogen pressure, and electrolyte

concentration were the parameters that were varied.

All electrodes were stored in the discharged state, open

circuited; in boiler plate pressure vessels, stacked on cores

that approximate flight hardware. They were separated by gas

screens. No mechanical pressure was exerted on the stack.

The vessels were evacuated and backfilled with hydrogen to the

desired pressure. The zero psig vessels were backfilled with

281
hydrogen then the pressure was released until the gauge read

zero. The vessels were then stored under the appropriate

conditions and monitored to ensure that the desired storage

conditions were maintained throughout the twenty-eight day

storage period.

After Storaq_e Procedures

The electrodes were removed from the vessels following twentyeight days of storage at the specified conditions. The

electrodes were placed in a solution of KOH at the proper

concentration. Open circuit voltages were measured vs Hg/HgO as

soon as possible following removal from storage. Impedance

measurements were made in the discharged state at the electrode

open circuit potential over the same frequency ranges as the

before storage measurements.

The electrodes were then set aside, flooded with the proper

concentration KOH, until the entire group of thirty electrodes

had completed storage and initial impedance evaluation at the

specified conditions. The after-storage capacity measurements

were made in the flooded state following the same procedures used

for the initial capacity measurements. Again five to six cycles

were run and the results averaged. The average before and after

capacities were compared in order to determine the percent change

in capacity for each electrode. Data from the groups of three

was also averaged for comparisOn purposes.

Following the after-storage capacity measurements, impedance was

measured after equilibration at 0.2 volts for 1 hour as in the

before capacity measurements. After storage measurements were

not made at 0.4 volts.

NICKEL-HYDROGEN CELL STORAGE TEST PROCEDURES

Cell Construction Parameters

Representative cells were built to evaluate state of charge on

storage in completed cells. Due to electrode availability and

configuration, only nickel electrodes from Hughes were used for

this portion of the study. Three stacks composed of three pairs

of electrodes per stack were constructed. The standard Air Force

recirculating design configuration was used. The hydrogen

electrode used was of the standard Air Force design. The

separator used for all three cells was a potassium

titanate/polyethylene blend prepared under a grant with Miami

University.

282
Before Storage Impedance Measurements

The cell impedances were measured following activation.

Impedances were measured at the cell open circuit voltage over

the range of 0.0007 Hz to 5.65 Hz.

Before Storage Capacity Determination Procedure

The formation procedure used on the cells was similar to that

used on the nickel electrodes. The cells were charged at C/10

for 16 hours and discharged at the C/4 rate. Capacity to 1.0

volt was measured. The formation cycles were repeated until a

stable capacity was achieved. Four cycles were run.

The pre-storage capacity was determined at the C/2 charge rate

with a 10% overcharge and a C/4 discharge to 1.0 and 0.1 volts.

Nine pre-storage capacity determination cycles were run.

The impedance measurements were repeated before storing the

cells. The measurements were made at open circuit following the

discharge to 0.i volts at C/10. They were made over the same

frequency range as the before formation measurements.

Storage Procedures

The storage procedures and conditions for the cells are

summarized below. Cell 1 was discharged at the C/4 rate to 1.0

volts then discharged at C/10 to 0.i volts. The cell was then

stored at room temperature in the open circuit state for twentyeight days. Cell 2 was discharged under conditions similar to

those used for cell 1 then stored at room temperature, shorted

through a 1 ohm resistor for twenty-eight days. Before storage,

cell 3 was charged at the C/2 rate with a 10% overcharge. The

cell was then trickle charged at C/200 for the twenty-eight day

storage period. In order to emphasize any pressure effects, all

cells were stored at room temperature and at 500 psi hydrogen.

After Storage Procedures

Following storage for twenty-eight days, the cells were charged

and discharged as in the pre-storage capacity measurements. The

cells were charged at C/2 with a 10% overcharge and discharged at

C/4 to 1.0 and 0.1 volts. A number of cycles were run and

capacity trends were monitored.

Impedance tests were repeated following the after-storage

capacity measurements.

283
RESULTSAND DISCUSSION

Nickel Electrode Storaqe Tests

Capacity Loss Measurements

The capacity loss suffered as a result of storage at the various

conditions is summarized in Figures 1 and 2. Figure I compares

capacity loss as a function of storage conditions for each of the

eiectrode manufaCturers. In general,_elec_0_es st0red in 31%

KOH, at 500 psi, and room temperature experienced the greatest

loss in capacity. This was consistent for all four

manufacturers. Storage at -20 deg C helped to minimize the

capacity loss even at hydrogen pressures of 500 psi. The low

temperature storage resulted in a .6 to 13% decrease in capacity

loss with other storage parameters remaining constant. Similar

reductionS, 4 to 13%, were aChieved by redu6ing_t_ _£orage

pressure from 500 psi to 50 psi. No additional reduction in

capacity loss was achieved by reducing the storage pressure to

ambient conditions. Electrolyte concentration appeared to have a

minimal impact on capacity loss with storage in the presence of

hydrogen. _

Figure 2 compares the capacity loss as a function of electrode

manufacturer for each of the storage conditions. Nickel

electrodes from Hughes exhibited the greatest capacity loss for

all storage conditions. This may be a result of the impregnation

method or the concentration of cobalt in the electrodes. Hughes

is the only one of the manufacturers whose electrodes were tested

that uses an alcoholic impregnation process; all others use

aqueous methods. The Hughes electrodes also have the highest

cobalt level of those tested (Table i). It has been postulated

that cobalt segregation in reducing environments may be

responsible for performance changes in nickel electrodes (ref.

7). Thus cobalt levels may effect the capacity loss

characteristics as well. In order to evaluate the effects of

cobalt concentration, electrodes impregnated by the same process

have been prepared with 5% and 10% cobalt levels for impedance

testing and capacity loss on storage evaluation.

In general, the capacity loss experienced by eleCtrodes from the

manufacturers other than Hughes was minimal; less than 5%. In

fact the Eagle Picher electrodes experienced a slight increase in

capacity measured after storage compared to that measured before

storage.

Impedance Measurements

Electrodes from each manufacturer exhibited characteristic

impedance behavior when measured in 31% KOH, at 0.2 volts vs

Hg/HgO as shown in Figure 3. However, when measured at 0.4 volts

vs Hg/HgO, the impedance spectra were similar for all four types

of electrodes and could not be differentiated. It appears that

284
the diffusion resistance increases abruptly as the electrode

voltage decreases (ref. 8). This in turn effects the kinetic

resistance term. The point at which, and the extent to which the

diffusion resistance increases appears to vary with manufacturing

process as well as previous cycle history.

The electrode model used for the impedance studies is shown in

Figure 4. The impedance values that were compared for the

various electrodes are for what is generally referred to as

kinetic resistance, R k. The kinetic resistance is a measure of

the resistance to electron transfer. The ohmic resistance is not

a valid parameter for the electrode tests as the measurements

were made in the flooded state and as such the ohmic resistance

becomes a measure of the resistance of the electrolyte to the

reference electrode. Since impedance was measured before storing

the electrodes, the "pre-storage' measurements on all electrodes,

with the exception of those tested in 26% KOH, were made at the

same conditions for all electrodes, i.e. room temperature and 31%

KOH. The initial, pre-storage kinetic resistance measurements

generally ranged between 0.5 and 2 m-ohms. The post-storage

impedance values are shown as a function of storage conditions

and electrode manufacturer in Figures 5 and 6. There was no

direct correlation in the changes in R k as a function of storage

conditions or electrode manufacturers. The impedance values seem

to correlate with the capacity loss observed within a group but

not between groups. Electrodes that experienced a capacity loss

after storage also exhibited changes to the impedance spectra

while those with no capacity loss did not (Figure 7).

The measurement of cell voltage upon removal from storage seemed

to correlate very well with the capacity loss measured in the

electrode tests. The Hughes electrodes stored at 500 psi and

room temperature lost 15% to 18% of their initial capacity. They

were also the only electrodes to have a negative voltage vs

Hg/HgO immediately following storage. In general, the lower the

electrode voltage upon removal from storage the greater the

capacity loss. Table 3 shows the electrode voltage vs Hg/HgO

upon removal from storage for each of the manufacturers at the

various storage conditions. The voltage upon removal is an

indicator of the amount of active material reduction that has

taken place in the nickel electrode. The impedance then follows

from the voltage change. In general, the lower the voltage, the

higher the electrode impedance.

Cell Tests

The capacity to 1.0 and 0.1 volts vs cycles for the cell storage

tests are shown in Figures 8-10. The results for the cell stored

discharged/open circuit and the cell stored in the shorted

condition are similar. In both cases the cells exhibited a

capacity loss of 17-18% when comparing the cycles immediately

prior to and after storage. This capacity loss was in the same

range as that seen for Hughes electrodes stored under similar

285
pressure and temperature conditions in the electrode portion of

this study. The electrode stored at 500 psi and room temperature

experienced an average 15% loss in capacity over the twenty-eight

day storage period. However, the capacity loss experienced by

the cells was recovered within eight cycles as shown in Figures 8

and 9. This does not agree with the electrode studies which

showed essentially no increase in capacity during the afterstorage cycling. The storage test was repeated on the same cells

and similar results were obtained the second time around.

The cell stored charged then trickle charged at C/200 for twentyeight days showed no degradation in capacity due to storage. A

plot of the capacity obtained Vs cycles for cell 3 is shown in

Figure i0. It should be noted that the difference between the

capacity to 1.0 volt and the capaci{y to 0.i volt increased

after the storage period. The capacity to 1.0 volt remained

constant while the total cell capacity increased.

The cell tests showed no indication of irrecoverable capacity

loss as a result of storage for any of the conditions tested.

Capacity recovery which occurs within twenty cycles of storage is

not considered a problem.

Comparison of impedance measurements made before and after

storage were inconclusivel ' There was a great deal of scatter in

the data and the absence of a reference electrode complicated the

measurements.

CONCLUSIONS AND RECOMMENDATIONS

Capacity loss appears to be related to the hydrogen pressure in

the vessel. Low temperatures, (-20 deg C) seem to negate the

effects of the high pressure. KOH concentration has minimal

effect on capacity loss. Nickel electrodes from Hughes showed

the highest capacity loss. This may be related to impregnation

process and/or cobalt levels in the electrodes. The storage

tests run to evaluate the effect of state-of-charge on storage

showed essentially no capacity loss for any of the conditions of

storage evaluated.

The capacity loss on storage study provided much valuable

information relating to the interactions of temperature, pressure

and electroly_e concentration with nickel electrodes from various

manufacturers. However, in some cases, conflicting resultS were

produced, specifically in the correlation of the electrode data

with the cell storage data. As mentioned above, areas requiring

further study have been identiffed and are being pursued.

Impedance work is continuing and as a greater understanding of

the data is achieved further correlations may be made and a

deeper understanding gained from the present study.

286

Z

=
ACKNOWLEDGEMENTS

The work described in this paper represents a group effort. The

following individuals were involved in completing the study:

Olga Gonzalez-Sanabria and John Smithrick developed the test

matrix used to evaluate electrodes from the various

manufacturers. John Smithrick coordinated efforts with the

manufacturers and obtained the electrodes necessary to perform

the work. Michelle Manzo planned, scheduled and coordinated the

overall test scheme, set up the flooded capacity measurement

apparatus, and monitored the storage of the electrodes. Jeffrey

Brewer, Doris Britton, and Olga Gonzalez-Sanabria were

responsible for running the flooded capacity test measurements

before and after electrode storage. Randall Gahn wired the data

acquisition system for capacity loss measurements, prepared

pressure vessels for both the cell tests and electrode storage

tests, and built and tested the cells for the state-of-charge

storage evaluation testing. Margaret Reid provided all of the

impedance measurements on both the electrodes and cells. Jo Ann

Charleston organized and plotted the results. Patricia O'Donnell

and Lawrence Thaller served as advisors to the effort.

i •

.

,

•

•

6

.

REFERENCES

McDermott, J.: Low-Earth-Orbit Testing of Nickel-Hydrogen

Cells. 21st IECEC, Advancing Toward Technology Breakout in

Energy Conversion, August 1986, pp. 1516-1520.

Vaidyanathan, H.; Dunlop, J.: Capacity Maintenance for

Intelsat VI Nickel-Hydrogen Cells. 21st IECEC, Advancing

Toward Technology Breakout in Energy Conversion, August 1986,

pp. 1516-1520.

Vaidyanathan, H.: Long-Term Storage of Nickel-Hydrogen Cells.

Space Electrochemical Research and Technology Conference,

NASA Lewis Research Center, Cleveland, Ohio, April 1987,

Proceedings, pp 41-48, NASA Conference Publication 2484, pp.

41-48.

Smithrick, J.: Effect of Storage and LEO Cycling on

Manufacturing Technology IPV Nickel-Hydrogen Cells.

New Frontiers, 22 IECEC, August 1987, pp.878-881.

Energy

Vaidyanathan, H; Earl, M.: Capacity and Pressure Variation of

Intelsat VI Nickel-Hydrogen Cells with Storage and Cycling,

1988 IECEC Proceedings, pp. 471-475.

Hartjen, G. L.; Applewhite, A. Z.; Hall, A. M.:Storage and

Testing of 220 Ampere-Hour Nickel-Hydrogen Battery Cells in

Low Earth Orbit, 1988 IECEC Proceedings, pp.363-366.

Zimmerman, A. H.: Nickel-Hydrogen Cell Positive Electrode

Studies: Cobalt Segregation in Reducing Environments, Report

287
•

SD-TR-87-64, The Aerospace Corporation, E1 Segundo, CA, May

1987.

Reid, M. A.: Impedance Studies of Ni-Cd and Ni-H 2 Cells Using

the Cell Case as a Reference Electrode, Space Electrochemical

Research and Technology Conference, NASA Lewis Research

Center, Cleveland, Ohio, April 1989.

288
Table 1

Nickel Electrode Characteristics

Characteristic

Plaque type

Plaque porosity (%)

Impregnation process

Loading level (g/cc void)

Cobalt level (mole %)

Thickness (mils)

Hughes

Dry Powder

8O

Alcoholic

1.60

i0

3O

EaglePicher

Slurry

8O

Aqueous

1.65

5

30

Whittaker

Yardney

Slurry

8O

Aqueous

1.68

5

35

Gates

Dry powder

8O

Aqueous

1.63

31 or 37

* Proprietary information

Table 2

Nickel Electrode Storage Conditions

H2 Pressure

(psig)

5OO

5O

0

5O0

5OO

Temperature

(deg C)

RT

RT

RT

RT

-20

KOH Concentration

31

31

31

26

31

Code

HPRT31

MPRT31

APRT31

HPRT26

HPLT31

289
Table 3

Nickel Electrode Voltages After Storage

Measured vs Hg/HgO)

Storage Condition

HPRT31

MPRT31

APRT31

HPLT31

HPRT26

Hughes

-0.765

0.186

0.212

0.199

-0.365

Electrode Manufacturer

Eagle

Picher

0.173

0.198

0.240

0.230

0.216

Whittaker

Yardney

0.032

0.191

0.216

0.221

0.073

Gates

0.170

0.200

0.244

0.233

0.216

290
18

16

14

Figure I

St,orage ConOli,Jons

n _o m ,,, ,, _ 0=

Capacity Loss vs Storage Conditions for Nickel

Electrodes from Various Manufacturers

ORIGINAL PAGE IS

OF POOR QUALITY

18

%6

%4

,0J

e.

o

u

¢p

in

In

o

4J

_4

u

12

10

8

6

0

-8

-•

-6

-II

m--,, mm=,,, m _, w =o, m .,_,

7£guro 2 Capacity Loss vs Nickel Electrode Manufacturer for

Various Storage Conditions

_ • =_.

291
Gates

11,4

|

t

@

v

II #'

e#

/

t !

|

@

w

i

o

! i • ' ' O.it n

Z' (ehu)

Eagle Picher

|

£

0

I

|

0

%,.

!

"1.,4

Hughes

i • | • . s

Z" (ehu)

,.s Whittoker-Yordney

0

e

e

• , j • i i I I I I I

2

2' (ll_) Z' (d_em)

| s

=,8

. , -7,

Figure 3 Representative Impedance Curves for Nickel

Electrodes from the Various Manufacturers

Impedance measured at 0.2 volts vs Hg/HgO

CDL

IRk Z w

RI_ OHMIC RESISTANCE

CDL DOUBLE LAYER CAPACITANCE

R k KINETIC RESISTANCE

Z w WARBURG IMPEDANCE

ri_:e 4 Electrode Model Used for Impedance of Nickel

Electrodes in Storage Studies

292

ORIGINALPAGEIS

OFPOORQUAUn'
&

M

90

80

_0

40

10

0

ORIGINAL PAGE IS

OF POOR QUALITY

Figure 5 After Storage Impedance (Rk) vs Storage Conditions

for Nickel Electrodes from Various Manufacturers

ql(3

Figure 6

Xloct.rodo lfanuCscturo¢

mm m,.. m ---, m .,_ m -_ m -_,

After Storage Impedance (Rk) vs Nickel Electrode

Manufacturer for Various Storage Conditions

293
i

£

o

!

Electrode Stored at Ambient Pressure, Room Temperature

2

_._ Before Storage

l ' J ,l I I I I

l S ( el'ms )

41

I

,-

0

v

!

After Storage

I i i J i • i

ee

Z" (d'ul)

II

,C

0

v

I

Electrode Stored at 500 psi, Room Temperature

Before Storage

l I l I I I I

I

,C

0

!

__After Storage

l • I I . ! l l

Z' Ides)

YLgure ? Impedance Spectra of Nickel Electrode Exhibiting

Capacity Loss After Storage (shown at top) Compared

to Impedance Spectra of Nickel Electrode That Did

Not Exhibit Capacity Loss (shown at bottom)

p

O

=

).

_q

u

4. |0

• II to t.O volt + J_ to O.i volt

Figure 8 Capacity vs Cycles for Nickel-Hydrogen Cell Stored

Discharged in the Open Circuit State

294

ORIGINAL PAGE IS

OF POORQUALITY
_.SO !

,in

i no

3#0-

3.80 -

3.70 --

o 3.60 -

3. SO -

,_. 3.40 -

-,_ 3. l0

3,10

U 3.00

_1.90

t. 80

Z. "tO

|, 60

|. UO

PrI-S_O_Ifl C?CII! /

• JU[ to t.O Tolt

CTcle N=mber

Figure 9 Capacity vs Cycles for Nickel-Hydrogen Cell Stored

Discharged, Shorted

4.20

4.10 -

4. DO

3. gO

3. (10

3. ?0

i 3.60

3_ SO

3, tO

3. $0

3. IO

_, 3. tO

$.m

I.N

1.80

8.U

3.N

P[o-S'r, orqtgl Cyclos . _

, _ , _ , _ , _ ,

CycZe giber

• kll to I. 0 _mlt 4- Ii tlo O. JL _lt

Figure 20 Capacity vs Cycles for Nickel-Hydrogen Cell Stored

on Trickle Charge at C/200

ORIGINAL PAGE IS

OF POOR QUALITY

295
i!i

=

.]
N90-20479

SODIUM-SULFUR BATTERY FLIGHT EXPERIMENT DEFINITION STUDY*

Rebecca Chang

Ford Aerospace Corporation

Space Systems Division

Palo Alto, California 94303

and

Robert Minck

Ford Aerospace Corporation

Aeronutronic Division

Newport Beach, California 92663

The need for high power space systems is anticipated within the

next ]0 to 20 years. Examples are the Space Station, co-orbiting

platforms, geostationary orbit (GEO) platforms, and space based

radar satellites and some classified high power missions. Sodium

sulfur (NaS) batteries have been identified as the most like]y

successor to nickel hydrogen (Ni-H z ) batteries for space

applications. One advantage of the NaS battery system is that the

useable specific energy is two to three times that of Ni-H 2

batteries. Th:is represents a significant launch cost. savings or

increased payload mass capabilities. NaS batteries support NASA

CAST's proposed Civil Space Technology Initiative goal of a factor

of two improvement in spacecraft power system performance, as well

as the proposed Spacecraft 2000 initiative.

The NaS battery operates at between 300 and 400°C, using liquid

sodium and sulfur/polysulfide electrodes and solid ceramic

electrolyte 1'2'3 The transport of the electrode materials to the

surface of the electrolyte is through wicking/capillary forces.

These critical transport functions must be demonstrated under

actual microgravity conditions before NaS batteries can be

confidently utilized in space.

Ford Aerospace Corporation, under contract to NASA Lewis

Research Center, is currently working on the NaS battery space

flight experiment definition study. The objective is to design the

experiment that will demonstrate operation of the NaS battery/eel]

in the space environment with particular emphasis on evaluation of

microgravity effects. Experimental payload definitions have been

completed and preliminary designs of the experiment have been

defined.

* The work at Ford Aerospace Corporation is sponsored by NASA

Lewis Research Center under contract NAS3-25355.

PRECEDING PAGE BLANK NOT FILMED

297
ORIGINAL PAGE fS

OF POOR QUALITY"

INTRODUCTION

'Ih_ NaS batter_ cell, pioneered by Ford Motor Co., is seen as

ohe ut" the most attractive battery systems for spacecraft power

sj.'slem._. The primary advantage of this system is its high energy

d,:.rlsity, ;;hioh can reach 150 WH/kg at C-rate disoharge at the cell

levt-1. The NaS battery offers a eombinat:ion of system level

adva!_t ages "

O Nas;s azLd volume effeet, iveness

U [|igh rou]l,l-trip _t'fi,zi_:ncy

O tloat rt-.j(rL-tion at high tenlperatllre

O Simi)i i(,.i ty of battery integrat ton

O COml;_-t itive acquisition costs

Nominal NaS c_] 1 Operating parameters are sho_:n and compared

wi ill typical Ni-Cd and Ni-H 2 cell data in Table 1 . Unusual,

pert't, rmanve f_atures for the NaS cells are tire high cell operating

_ult, ages, high roundtrip ,_fficieJ_cy and high operating temperature

t:hich mlnimizes the radiator area and cost. Potentiai cost savings

are i l [ustrar.ed in Table 2. It is obvious that the battery mass

savings are t. ransl ated into spadeeraft launch cost savings or

p:_yload capability inorease. [!sing Ni-H 2 systems as a reference, it

('an ._ave in laun,,h cost from I .1 to 1.7M$/gW for the GEO missions

ai_d 0.2 to 0.4M$/KW for the LEO missions, For exampie, a 2.5 KW

aorospa(:e Na._ buttery concept design which weighs 31.8Kg while a

0.5 l(W Ni-H z battery _,eighs 30.5 I{g. o

Tl,e Nag batter}" is a particularly good energy storage system

for high po_4er satellite missions. A parametric study comparing

the operat, ional storage candidates for a 75 KW satellite mission

has been performed by Ford Aerospace and the results are shown in

Table 34. Tire NaS battery looks the most favorable with respect to

the other storage syst.ems .....

The NaB battery ceils operate at between 300 and 400 ° C. The

liquid electrode materials rely on the wicking actions for

tr_msport to and from the e].ectrolyte interface to support routine

c_c-l, ic _ operation. Current deveiopmenta] space cells have been

d,-signed to h_,ve adequate wicking, but this critical function must

be demolisi;r_t_--d under - actual orbYtal microgravity conditions before

NaS bat teries can be confidently applied in space. The primary

object ire of the Ford Aerospace study is to define the specific

e:.:perimehtal design required for successful demonstration of NaS

L:_l:tery/celIs under space environments.

:Z:/; _2IE :3_ ZT:

CELL OPERATIONS AND SELECTIONS

N,_S cells as sho,:n ir, Figure 1 utilize a solid ceramic

etc,_._.rolyt,- kno_:n aa _"-;.lumina to separate anode (sodium} and

c_:,-,.h,,d,r I >.ul fur/s_di t_,, pol ysu] fide_ ) reactants and to provide a

=

298

...... :: : 1:2-::_2 2_ :2: :::

= = .

=

_=:

_=

i
ORIGI_iAL PAC.;_ tS

OF POOR QUALITy

conductive path for sodium ions during operation. Tile transport

of tile elect.rode materials to the surface of the electrolyte is

through _;icking/capillary forces. A metal protection tube limits

the flow ,_f' sodium into the reaction zone, thus providing for safe

operation of the ceil. Since only small amounts of the sodium are

available for reaction, a quick, explosive reaction is not possible

to occur even in the case of a broken electro]yte tube. The sulfur

is contained in the outer part of the cell, imbedded in a carbon

fiber mat. The carbon fiber mat improves the electrical conduction

path for the sulfur electrode.

At the _.ell operating temperature, the reactants are at liquid

state. During discharge of the cell, the sodium splits into

positive ions, which are transported across the electrolyte, and

negative electrons, uhieh are carried up by the current coliector

to the _xternal circuit. Once the sodium ions are transported

across the ceramic electrolyte, they join with the sulfur to form

a sodium-polysulfide (Na2Ss). As tile cell is further discharged,

the Na2S 5 reacts _-ith sodium to form other sodium polysulfides such

as Na2S _ and Na.,S_. Charging the celI results in the opposite

reactions taking place with the sodium-pol ysul fides being

dismantled into sodium ions and sulfur. The sodium ions are then

transported back across the electrolyte to the sodium eiectrode.

The proposed flight experiment cells which are rated .10AH are

derived from the Ford Aerospace baseline SATBAT-2 ceils. The

physical characteristics of the cells are summarized as follows and

typical 40AH cells are sho_'n in Figure 2.

Cell Capacity (Ah): 40.0

Outer diameter {cm): 3.6

Cell length (cm): 23.0

Cell Mass (g): 519.0

FLIGHT EXPERIMENT JUSTIFICATIONS AND OBJECTIVES

NaS cells differ from most batteries in that they contain

molten anodic and cathodic reactants. Ability of the NaS cells to

efficiently charge and discharge is critically dependent upon the

favorable spatial distribution of the fluid reactants with respect

to their interface with the solid beta"-alumina ceramic

electrolyte. Fluid motion and reactant morphology are particularly

critical for NaS cell operation since the sulfur/sodium-polysulfide

catholyte reactants form three immiscible phases during recharge,

two of which are non-conductive: molten sulfur, void volume, and

Na2S 5 . Should an unfavorable distribution occur, a blocking layer

could develop and restrict cell operation.

Terrestrial cell designs incorporate preferential _etting,

capillarity and/or graded resistance uithin the sulfur electrode

to maintain efficient operation. Such ceils have been de_eloped

to operate effectively at high rate and to yield high specifi,,

299
ORIGINAL PAGE IS

OF POOR QUALITY

energy and specific power. However, the characteristics of NaS

cells under microgravity conditions are unknown because of the

uncertain role of gravity in controlling the distribution and

conso]idation of the separate reactant phases These effects can

only be determined by testing in orbit with 3-axis stabilized

spacecraft because i t is not. possible to simulate low gravity

cc)nditions toy sufficiently long times to represent cycling or the

large equilibration time constraints associated with the sluggish

flui,ts.

The main purpose of the NaS battery/cell flight experiment is

to validate the application of NaS battery technology to space

po_er applications. The principal objective of the flight

experiment definition study is to design an experiment, that will

demonstrate operation of the battery/cells under space environments

_;ith particular emphasis on evaluation of microgravity effects

_hich are specially critics] for 3-axis stabilized spacecrafts.

These call be categorized in the foli0w{ngs:

T. To evaluate cell charge and discharge characteristics as

affected by the fluid reactant distributions.

II. To cletermJne reactant distributions under microgravity

conditions.

ITl. To understand current and thermal distributions within

the eel ] s.

TV. T_, evaluate freeze/thaw effects.

V. To evaluaze multi-cell early LEO (Low Earth Orbit) cycle

life.

The approach is to select only those tests that are critical

and expected to differ under microgravity conditions. Spin

stabilized satellite batteries can be tested on the ground, thus

a test to simulate this application in space is not needed. Warm

launch has an advantage that the cells in space will require less

energy and time to heat up from the warm launch temperature to its

final[ ol)erat, ing temperature. Ho_ever, the demonstration of the

cell survivabillfy=£o the launch environments can be confirmed on

the ground by test fo] ]owed by electrical cyclic testings.

Therefore, the warm launch test need not be included in the

exper imerit.

DESCRIPTION OF PROPOSED EXPERIMENTS

Five tests have been iden{ified in the NaS battery flight

experimel,t definition study to meet the above objectives. Each of

these tests will be summarized in the following sections.

I. CELL CHARACTERIZATION TEST

II. REACTANT DISTRIBUTION TEST

llI. CURRENT/TEMPERATURE DISTRIBUTION TEST

IV. FREEZE/THAW TEST

V. MI!I.TI-CI_LI, LEO CYCLE TEST

300
P ORIGINAL PAGI _

OF' POOR QUALiTy

The test cells will be Ford Aerospace baseline 40AH cells

except Test III cells which will be specially instrumented with

extra thermocoup]es and voltage probes.

Our general testing philosophy is to have most of the space

cells evaluated under varim]s conditions on the ground prior to

launch, in the space, and after cells are brought back to earth to

provide before-during-after data comparison. Additional control

cells will be assigned to an identical ground test to provide

additional comparative data base.

I. CELL CHARACTERIZATION TEST

NaS cells will be evaluated under various test conditions in

space under two temperature ranges, 275°C and 350°C. Once orbit

has been stabilized, the cells will be heated to operating

temperatures and the test sequence performed. The charge/discharge

cycle variables have been established to cover a sufficiently wide

range of parameters to address a majority of anticipated space

applications. Pulse loads will be imposed during the discharge.

Each cell will be instrumented with power leads, voltage sense

leads and a thermocouple pair, all which will penetrate the thermal

insulation barrier for connection to ambient, temperature connector

panels. During the charge/discharge cycle, each cell voltage will

be periodically measured.

IT. REACTANT DISTRIBUTION TEST

After completion of Test I, the cells will be discharged into

their l-phase region and placed on open circuit to accurately

determine their state-of-charge. Each state of charge of the cell

will be adjusted to a composition ranging from NazS _ to Na2S 5.

III. CURRENT/TEMPERATURE DISTRIBUTION TEST

This test will determine the spatial variations of electrode

reactions and current densities within the sulfur electrode

continuously throughout the electrical cycles while being operated

under microgravity conditions. This information will support

analysis and assessment of test results of Test I and If, and will

provide direction of possible modification of the sulfur electrode

design should that prove necessary.

The distribution of internal current densities and electrode

reactions will be measured continuously during the electrical

cycles by measuring its small ohmic losses. Specially instrumented

cells will be fabricated to incorporate a multitude of voltage

probes on the positive current collector. In addition to the

external thermocouples located at the top, middle, and bottom of

the cell container, internal thermocouples on the inside

electrolyte tube wall will also be installed at respective

locations to measure the temperature gradients. Three sets of

voltage probes shall be used to provide information about

circumferential uniformity.

301
ORIGINAL PACE K"

POOR QUAL!_ _

%

[V. FREEZE/TII.XW TEST

Cells _,:i 11 be evaluated For the fr_eze/thaw cycles. Cell._

_;il. 1 he _nitial]y heated up t.,_ operating temperature and evaluated

electrical [Y, tel lowed by several freeze/thaw cycles. On the last

day nf the F] i_ht experiment, cells _ill be again evaluated

e.l_ctrical ].v t-o determine the effects <_f th_ freeze/thaw llnder

miarogpavity conditians. The ce]l.q _'ill be cooled for reentry.

V. HUI,TT-MF.r,L I_EC_ CYCI,E TEST

._aS ee]]._ _.'ill he sorio_ connected and evaluated in the 90

minllfe c, rbit rogimr,, Cells _;il] be cycled at a high rate di._charge

f'¢_r _ minister and Fnl]o_..-ed by a charge for 54 minutes. It i.s

almn.qt impo_.qihle to ev_{luate t, he life cycle of the NaS cells Jn

s_,ch _hn,'t e-cperimant duratian. However, the proposed LEC) cycle

_;i] ] pr.nvide an early ] Ire Pea]_ time LEO operational experience and

demon.qirnt_ the multi-cell operation. .Tnformati_n obtained wi] ]

he definitely va]lmt_le for fu]_l size battery _ca] ing-up.

SYSTEH iHPI,EHENTATION AND CARRIER OPTIONS

•\ yap ot.y of payload carriers, shown in Figure 3, including

,_-Ii d le,'k Ho,q_ la r l,ocl;e_-, Get Avaay Special , Hitchhiker G, and

Hi f_hhikAV H h._ve been revia_-'ed s. Concerns such as electrical

pc, w_r, ma._._ Joad/._t.r_cture, and mounting Orientation have been

c,_nsid_P_d as Factors for carrier ._.IectJons. From the preliminary

d,_.qi_n of the Civa te._ts, the Lois[ _eight of the enclosure._,

batteries, and .qupimrt equipments is estimated to be a total of

l.o._ than lf)O lbs. The experiment _ill occupy a space

alU_rc, ximat_l.x/ 3.O'xR.O'xl .5'

SUMMARY AND CONCLUSION

NaS batteries have. been identified as the most likely

.q_cces._nr to _pace Ni-H_ or Ni-Cd batteries, primarily doe to a

m._._s r_duefion by a factor 222 0Vet Ni-H_ and by a factor of" 4 over

Xi'Cd; This yields me.jar lai_nch cost red{rations or payload mass

impvovoment ._ . ,kiss batteri es support NASA CAST ' s proposed Civil

Space T,_.chn_lngy Tnit_ative ._nal of a Factor of two improvement in

.qpaceceaft 2000 initiative, Sine._ Ni-H a and Ni-Cd batteries have

been .qD._C,e Fl i._ht prnven, j.t is e_sential to have the flight

_xpevim_nt fo estnbl_sh a national space technology base to

demon'+it:aLe= = t_he _perat i on ..... of the NaS battery for space

a I-'P I i e n t i c_n ._,

302
ACI4_JOwI.,EDGEHENT

ORIGINAL PAGE t8

OF POOR QU_.LI'F_

The authors wi.qh to thank our cnl _I_a_,,_s Far t e,qhni eel

assistance And the ,_uppoPf c.f thi_ ,_rork From NA.qA l,_,cis I_esearrb

Center under con.,ra'-o.i No. NAS3-25255.

REVERE_CES

1. R.W._linck & C.R.Ha]bach, "Characteristics mf Sodium-Sulfur Cel]_

for Divorae Applications", Prooeeding of the 171h IECEC, 19R2,

P.557

2. H.L.Haskins, H.L.HcC]anahan, and R.W.Hinck, "'Sodium-ElJlfur Cells

for High-Pnwer Spacecraft Batteri_", Prncoeding of t h_ lRth IIECEC,

1983. P.1465

3. D.H. Allen, "sodium-Sulfur SatelJife Batteri_: Ceil Te_t

Results and Development Plans", Pcoceeding_ of l he 19th IECEC,

1984, P.163

4. C.W.Koehlor & A.Z.Appleuhit.e, and A.H.Hall & P.G.Pltg_e].[,

"Bipolar Nic'kel-Hydrogen Battery Development", Proceodjn_ af th_

20th IECEC, 19R5

5. NASA Marshall Space Flight Center, "STg Tnvestigs%_rz G_ide"

303
ORIGINAL P,_.GE IS

OF POOR QUALITY

Table !.

%

SPACE BATTERY OPERATING PARAMETER COMPARISON

I Battery

I System

] -

I

i Ni/Cd

I

I

i_

I

INi/H 2

(IPV)

NI/H 2

(bipolar)

NalS

I Average Cell IRoundtrip I Discharge I Operating ITypical

Condltion I Voltage !V/_____IEfficieneyIDisslpation}Temperaturel DOD

(%) [ (kWt/kWe) [C3______J______

GEO

LEO

GEO

LEO

GEO

LEO

[Char_ejDischarge[

I

1.43 1.22 I

I

1.45 1.23 I

I

I .45

1.47

I .45

1.46

2.10

2.15

I

so I

I

8o I

1

I I

1.25 I 8o I

I I

1.27 I so I

I .I

I I

1.26 I 82 I

1 I

1.28 I 82 I

! 1

I I

1.83 I 87 I

I I

1.so I 8_ I

I J

GEO

LEO

0.18 0-20

O.17 0-20

0.20 0-20

0.18 0--20

O. 19 0-20

0.27 0-20

0.18 300-400

0.21 300-400

I

I 45-6o

I

115-22

I

I 68-eo

I

1 28-4o

[___

I

I 68-80

I

I 28-40

I

I 50-80

I

I 35-50

I

Table 2.

[ Battery I

Sodium/Sulfur Batteries Offer Spacecraft Mass Savings

Mass i Volume I Cost- vs Ni/H2 ILife I

[__stem|Condition

I I

l.i/Cd I GEO

I I LEO

I L

I I

INi/H2 I GEo

I(IPV) i LEO

I [

I I

INi/H2 i cz0

Perfotnnance l Performance J2£o_d

I I

55-90 kg/kW [30-55 1/kW" [ I.I

70-140 kg/kW 140-85 I/kW [ 1.6

i i

I I

35-60k_/kW 160-120 1/kW I 1

35-70 kg/kW ]60-150 I/kW [ 1

I I

I I

35-60 k_/kW ]25-_5 Z/kW I 1.2

125-55 1/kW I 1.2

I I

Launch

+ 1.0-1.5 M$1kW

+ 0.3-0.5 M$1kW

0

[Years [

I

5-7 I

5-5 I

I

I

lO-15+1

5-s I

I

I

10-15+i

](Bipolar)[ LEO 35-70 kg/kW

I I

I I

[Na/S [ GEO 13-25 kg/kW

[ [ LEO 11-20 kg/kW

I k

a) requires demonstration

I I

14o-12o I/kW I 1.2

130-85 1/kW I 1.2

1 1

- 1.1-1.7 M$1kW

.- 0.2-0.4 H$/kW

5-8

5-(a)

1-(a)

=

=

|

=

i

m

F.

i

=

=

=

=

304
ORIGINAL PAGE _',_

OF POOR QUALITY

TABLE 3. ENERGY STORAGE SYSTEM OPTIONS COMPARISON

ENERGY STORAGE SUBSYSTEM OPTIONS

PARAMETER

ROUND TRIP EFFICIENCY (%)

DEPTH-OF-DISCHARGE (%)°

MASS (kg)

VOLUME (m 3)

WASTE HEAT/CYCLE (kW-h)

ECLIPSE HEAT REJECTION (kW)

TEMPERATURE {°C)

CHARGE POWER REQUIREMENT (kW)

"IN ALL CASES. EXCEPT Ni-Cd CONlrROLLED BY

FOLLOWING A FULL POWER ECLIPSE

°EXCLUSIVE OF CONTAINMENI

RFC RFC I OlPOL IPV FLYH2"O2 H2-O2 I NJ't42 N_-)I2 N_-Cd Na S WHEEL

3j 3T---!-a-T-- -- a--E---

21oo _5oo 3_oo 4_oo 12200 _5oo 42oo_

9.3 9.9 2.5 6.6 [ 5.2 1.9 a.0"

42 _4 15 17 I 17 11 I1

i

57 )5 17 18 I 17 17" 9

80 80 I0 10 L _o 350 35

....... I .....

98 )(2 _74 .--_,76 I 76., 66 i 66 .

ONE ORBIT HALF POWEr1 REQUhqEMENT

FI(_O_-I. SCHEMATIC OF NaS CELL

i

METAL CONTAINER--_

PROTECTION TUBE_ ">__)__._

SEALS -'-.,_

SODIUM._

METAL CONTAINER

BACKING ELECTRODE _

I__Na_+Na++e - _ Na+S--.,,. NazS

S+2er-'_'-S " I (2.08Volls) 5

F POLYSULFIDE

OF SODIUM LIMITED FOR

SAFSTY

(+)

305
ORIGINAL PAGE

BLACK AND WHITE PHOTDGRAPri

FI(_JRE 2. FORD AEROSPACE NaS CELLS

!

FIC_'RE 3. PAYDOAD CARRIER OPTIONS

_pt-_waV

Src_i.ls

MDM Pallet

"0 c ¢'<( ,_

3O6

ORIGINAL PAGE IS

OF POOR QUALITY

I

Lockers --t _i I _ _ " :<:-, _:_we_,_

=
ORIGINAL PAGE IS

OF POOR QUALITY

N90-20480

ADVANCED HIGH-TEMPERATURE BATTERIES

P.A. Nelson

Argonne National Laboratory

Argonne, Illinois 60439

Recent results for Li-AI/FeS 2 cells and a bipolar battery design have shown

the possibility of achieving high specific energy (210 Uh/kg) and high specific

power (239 W/kg) at the cell level for an electric vehicle application. Outstanding performance is also projected for sodium/metal chloride cells havin_

large electrolyte areas and thin positive electrodes.

INTRODUCTION

Work has been under way for about two decades on high-temperature battmrir_

having lithium or sodium negative electrodes. These efforts have met with -om_

success, but the original promise of very high specific energy and power has not

been achieved for practical battery systems. This paper discusses some rmccn1" -%"

approaches to achieving high performance for lithium/FeS 2 cells and _odlum/m=tal

chloride cells.

A comparison of the voltages and theoretical specific energies for several

high-temperature cell couples is shown in Table I. The fraction of the

theoretical specific energy that can be achieved for practical cells differ_

markedly among the couples shown in Table I. The important factors 5n determinin;

that fraction are the voltages of the cells and the densities of the reactants Pn _j

the discharge products. The densities of the materials have a m, rpr_ing]y

important effect on the achievable specific energy, as is discussed in mor_ deta:"

below.

LITHIUM/DISULFIDE BATTERIES

The main problems for the development of successful Li-AI/Fe_ cell_ have

been (I) instability of the FeS 2 electrode, which has resulted in rapidly

declining capacity, (2) the lack of an internal mechanism for accommodmt_ng

overcharge of a cell, thus requiring the use of external charge control on each

individual cell, and (3) the lack of a suitable curremt collector for the positi ....

electrode other than expensive molybdenum sheet material. Much prnKrcsm has b_r_

made at ANL in solving the first two problems, and _ new appromch to cell d_sigD

may result in a solution for the materials problem.

The disulfide electrode discharges along two major voltage plateau_, ar _h.-,:

in Fig. 1 (ref. 1). It has been found practical to limit the ammmt of lithi-m _

the negative electrode to stop the discharge at the end of the first major

plateau, with a final discharge product of Li2FeS 2. Although, this l_mits tbe

theoretical specific energy to about 475 _h/kg (second row, Table I. ba_ed on

Li-50 atom % A1), the positive electrode can be fabricated as a much den_er

electrode without provision for the large volume of Li2S produced in ODCra[tnq C _

the lower voltage plateau.

307
The early problem on the stability of the FeS2 electrode involved the

disassociation of FeS2 at high voltages on charge with the formation of

polysulfide ions, which dissolve in the electrolyte and then result in deposition

of lithium sulfide in the separator region.

ThomasD. Kaun at Argonne National Laboratory studied this stability problem

and also sought to achieve higher energy and power by the incorporation of dense

FeS 2 electrodes operated only on the upper voltage plateau (ref. 2,3). He found

that the use of dense FeS 2 electrodes not only improves the specific energy but

also improves the electrode conductivity and results in higher power. In most of

the earlier work, the electrolyte was the LiC1-KC1 eutectic (m.p., 352°C), which

required operation of the cell at about 450°C because of polarization of the

electrolyte with high composition gradients at high operating rates. Kaun

introduced the use of an LiC1-LiBr-KBr electrolyte having a lower melting point

and a broader liquidus range than those Of the LiC1-KC1 electrolyte, permitting

operation at 400°C. The combination of these changes improved the electrode

utilization by 50% and doubled the power capability of the electrode at 80% depthof-discharge (Fig. 2). These changes also resulted in a dramatic improvement in

cycle life capability, as shown in Fig. 3.

Another key improvement has been the development of a chemical overcharge

protection mechanism (ref. 4). By providing an excess of lithium relative to the

aluminum or other alloy elements, a high lithium activity at the negative

electrode is reached at the end of charge. This results in an increase in the

dissolved lithium concentration in the electrolyte as the cell approaches full

charge and, thus, about a twentyfold increase in the rate of self-discharge for

the cell. By maintaining a low trickle-charge rate for several hours at the end

of the charge, cells that are only partially charged have an opportunity to come

to full charge, while fully charged cells tolerate the overcharging by the

increase in the self-discharge rate.

Recently work has begun on bipolar Li-A1/FeS 2 cells. The bipolar

configuration reduces the amount of nonactive material required in the cell, and

it provides a simpler design for the use of a coated material to reduce the cost

of the disulfide electrode current collector. A schematic of the experimental

cells now under development is shown in Fig. 4.

Researchers at ANL have used the results of cell tests to project the

performance of Li-A1/FeS and Li-A1/FeS 2 cells in both the bicell and bipolar

configuration. The results are summarized in Table II. The capacity and power at

the end of discharge (EOD) of the four types of lithium/sulfide cells in Table II

were selected to be appropriate for the Eaton DSEP van being developed by Chrysler

and Eaton in a program for the Electric and Hybrid Propulsion Division of the

Department of Energy. The specific energy of the multlplate monosulfide cell is

essentially that which has been already achieved in the laboratory, but the

specific power assumes about a 50% improvement, which is expected to be achieved

with thinner electrodes. The bipolar monosulfide calculations are based on the

electrode performances used for the multiplate cells, with the weight of the cell

calculated by addfng the weight of individuai-parts. The bipolar monosulfide cell

would be approxlmately 7-1/2 in. in diameter and have electrodes 7-in. in

diameter. The performance of the disulfide bfceil is based on resuits achieved in

the laboratory for smaller cells, with the weight calculated from the individual

components. The calculations for the disulfide bipolar cell were based on the

monosulfide bipolar calculations, with substitution of parts as necessary for the

disulfide condlti0hs.

308
Table III gives the projected performance for batteries constructed for the

Eaton DSEP van. These results are based on the projected cell performances of

Table II. The initial range of the vehicle with a new battery would be 125 miles.

At the end of battery life it would be i00 miles. The energy storage required of

the battery was adjusted for the total weight of the vehicle, including the

battery, to achieve the desired vehicle range (Table III). It should be noted

that the vehicle weight is a function of the battery weight, being 1649 kg plus

the battery weight. The power at the end of discharge and at the end of battery

life is sufficient for accelerating the vehicle to 50 mph in 20 s. With new

batteries, for which the powers at the end of battery discharge are given in the

table, the acceleration would be approximately 20% more rapid than at the end of

life.

The number of parallel strings in the batteries varies from one for the

monosulfide multiplate to three for the bipolar batteries. For the bipolar

batteries, each string would be housed in a separate insulating container of

approximately lO-in. O.D.

Tests of very thin Li-AI/FeS 2 cells are appropriate for pulsed power

applications. Figure 5 shows the apparatus used by Laszlo Redey (ref. 5) for a

pulsed discharge measurement, and Fig. 6 shows the results that were achieved in a

discharge of approximately one millisecond. The calculated results show a

specific power of 82.7 kW/kg and a power-to-energy ratio of 574 W/Wh (Table IV).

SODIUM/METAL CHLORIDE BATTERIES

Sodium/sulfur batteries and sodium/metal chloride batteries have demonstrated

good reliability and long cycle life. For applications where very high power is

desired, new electrolyte configurations would be required.

The sodium/metal chloride cells are related to sodium/sulfur cells in that

they both use molten sodium negative electrodes and a _"-alumina solid

electrolyte. The sodium/metal chloride system was invented in South Africa and is

now being developed under the direction of Anglo America of South Africa. The

principal organizations involved in this development effort include Harwell and

Beta Research and Development Ltd. of England and Zebra Power Systems of South

Africa.

The sodium/metal chloride cells, unlike the sodium/sulfur cells, utilize a

secondary liquid NaC1-AIC13 electrolyte in the positive electrode in conjunction

with an active material of metal chloride such as FeCI 2 or NiC12 (ref. 6-16).

Because the NaC1 and the metal chloride are insoluble in NaC1-AIC14 at the

operating conditions, the NaC1-AIC13 electrolyte acts only to transfer sodium ions

within the positive electrode. The cells can be operated at temperatures down to

160°C (the melting point of the liquid electrolyte), but are typically operated in

present development work at 250°C to improve the kinetics of the positive

electrodes.

The voltages and theoretical specific energies are given in Table I, and the

overall reactions for the Na/FeC12 and Na/NiC12 cells are shown below:

2Na + FeCI 2 _ 2NaCI + Fe

2Na + NiCI 2 _ 2NaCl + Ni

309
ORIGINAL PAGE iS

OF POOR QUALITY.

_:_ _,_Ll_uctant feature of the metal chloride cells is that they are commonly

fabricated in the uncharged state. For the Na/FeC12 cell, iron powder is mixed

_xth NdCi po_der and pressed Into a positive electrode. The NaC1-AICl_

._iecL_ui¥_e is added to the top of the positive electrode material as a dry

....duL, a-d the cell is charged after heating to the operating temperature.

_i, _,,,porcanc feature of fabricating the Na/HC12 cells in the uncharged state is

that the reactants are relatively inexpensive and easily purified. The iron

j,uuder required for the Na/FeC12 cell is considerably less expensive than the

.,_ckei powder required for the Na/NICi 2 cell, an obvious cost advantage, but the

!_o/[liCi2 cell has better potential performance.

Host of the work on Na/FeCI 2 and Na/NiCI 2 cells described in the published

•iterature (ref. 6,7) is on cells of i00 Ah or even larger capacity. Batteries of

_uch cells are already being tested in electric vehicles and appear to be rugged

.,-d reliable. This reliability results from connecting the cells into a single

.eri_u-connected string. Unlike Na/S cells, Na/HC12 cells usually fail In the

.2v_=d-circuic co,dition. Large Na/HC12 cells can be connected in series, with

_he iuss of a cell resulting only in the loss of voltage associated with that cell

,=_d =_sentially no loss of ampere-hour capacity for the entire battery.

The consequence of fabricating such large cells with a single beta-alumina

;=c_rolyte tube is that the cells have only moderate power and large voltage

,._ wlien discharged at the three-hour rate. Also, to achieve even moderate

,:,scharge rates for the thick positive electrodes requires a high volume fraction

.,_ moiten salt electrolyte. This adds to the weight and, coupled with the large

t_duction in voltage from that of an open circuit, results in achieving a low

reaction of the theoretical specific energy. As an example, Na/NiC12 battery

c+lis of lO0-Ah capacity were reported to achieve a specific energy of about 109

_h/kg al,d have a cell resistance of 9 m_ (ref. 8). Uhen these cells were operated

....a 66-cell battery, the energy was reduced to 88 Wh/kg and 95 Wh/L, which are

....su than the desired values for an advanced battery. However, the small

..,Juction in specific energy of the battery compared to that of the cells is

,:_,:ou_aging and results from the high denslty and high packing density of the

•:_l]a. The life of both the cells and the batteries is reported to be excellent

,, ±000 cycles).

An important feature of the Na/HCl_ cell is that, in the positive electrode,

_i_ HCf z material is in thermodynamic equilibrium with the metallic constituent

_iron or nickel) and does not appear to attack current collector materials made of

:t,e sa_Jiemetal. In contrast, polysulfides in the positive electrode of

s_dium/sulfur cells tend to attack most metals and alloys except, perhaps,

chromium and molybdenum, which are quite useful as coatings.

Ue believe that much higher specific energies and powers than those a,hleved

:n the studies reported in the literature are possible with sodium/metal chloride

__lis. one indication of this is the high utilizations reported for Na/FeCI 2

,:ells at low current densities with relatively thick electrodes (ref. 7).

The sodium/metal chloride cell appears to have several unique

chai_acteristics. These include: (1) the cell failures are usually in the short-

= _'_i_cuit ,_ode, which provides very good battery reliability, (2) the reaction of

'_,u,. and eisner FeCl_ or NiCI_ following the fracture of an electrolyte is mild

_,h o,ly a slight temperature rise, a good safety feature, and (3) high active

:,-_=_ial densities and high packing density for the cells result in a high

310
calculated volumetric energy density for the battery and a lightweight, compact

battery insulating case.

To take advantage of these good features, the low power of the present cell

design must be circumvented by increasing the greater electrolyte surface area,

which could make possible the use of thinner positive electrodes.

One approach to designing high-performance sodium/metal chloride batteries is

to design very small cells with small-diameter, thin-walled electrolyte tubes.

Calculations at ANL indicate that for this approach the cells would have to be

less than 5 Ah in capacity for the electric vehicle application and, thus, may

result in high-cost cell fabrication and complex cell interconnection.

A more promising approach is to develop an electrolyte structure consisting

of many long, small-diameter tubes connected to a single header. The sodium could

be located either inside or outside the tubes. If the sodium is located inside

the tubes, the segregation of the sodium into separate compartments is a good

safety feature, and the heat which is developed in the positive electrode on

discharge would be more easily transmitted to the wall of the cell from outside of

the tubes. Whether the sodium is located inside or outside the tubes, each tube

may be provided with an individual current collector wire, which can be sealed

into the individual tubes. With this design approach, the header may be

constructed of a metal that is compatible with the positive electrodes (nickel,

for instance). Individual tubes could be sealed into the header with glass. It

may be possible to retain some of the flexibility provided by the metal header and

permit some movement between the individual electrolyte tubes, thus avoiding high

stresses and cracking that might result for a rigid header structure.

Another approach is to develop a flat-plate electrolyte structure such as

that shown in Fig. 7. Individual current collector wires could be sealed into the

square pockets, which are probably most appropriate for containing the sodium

electrode material.

A comparison of the characteristics of the three types of electrolyte

configurations (single large tube, multiple tube, and flat-plate compartmented) is

shown in Table V. A design cell capacity for each of these configurations is

shown on the top row of the table. The I00 Ah capacity for the single large tube

is typical of that now used in sodium/metal chloride cells. Cell capacities of

200 Ah are assumed for a multiple-tube cell and a flat-plate cell having

compartmented electrolytes. To achieve the assumed capacities for the respective

cells, I00 tubes are needed for a multiple tube cell, and 6 electrolyte structures

are needed for the flat-plate compartmented cell.

These structures in the assumed sizes provide considerably more electrolyte

area than that in the single large tube cell. The larger electrolyte areas

coupled with the thinner wall thicknesses of the electrolytes result in much lower

voltage losses through the electrolytes, as shown in the table. The electrolyte

dimensions selected for Table V are only representative of high surface area

designs and could be adjusted to achieve a wide range of power-to-energy ratio.

The voltage losses that develop in the positive electrode are also alleviated

by the use of electrolytes having high surface areas. In presently operated

sodium/metal chloride cells, the loading density of the positive electrude is

frequently about 0.25 Ah/cm 3. This would result in a radial width for an annular

positive electrode of 10.5 mm. At the high current densities required for the

311
small electrolyte area, the voltage loss in this electrode would be very high. At

the lower current densities for the high surface area designs, higher loading

densities should be possible, and a value of 0.35 Ah/cm 3 was assumed in Table V.

For cells with this higher loading density and higher electrolyte surface area,

the thicknesses of the positive electrodes are much lower than that for the

present single tube cell design. Thus, for these cells, the voltage loss in the

positive electrode would also be low even though the fraction of liquid

electrolyte in the electrode is lower that that for most of the cells reported to

date.

Comparing the two types of high-surface-area electrolyte configurations

considered in Table V, indicates that each has advantages. The multiple tube

configuration probably requires less new development effort for electrolyte

fabrication. Also, the tubular electrolytes would be more easily sealed into a

header than the flat-plate electrolytes. The flat-plate compartmented design

provides a more advantageous shape for the positive electrode than the multiple

tube configuration. For the flat-plate cell, the positive electrodes would

consist merely of flat plates pressed against a positive current collector sheet.

Provision for the current collector and the loading of the positive electrode in

the multiple tube cell would be complex, but clever designs may solve these

problems.

CONCLUSION

Much improvement is possible in the performance of advanced batteries over

that which has been achieved thus far. Future efforts in development of both

lithium and sodium anode high-temperature batteries are expected to result in

batteries that meet the goals that were originally set when development efforts

were first initiated: specific energy of about 220 Wh/kg, specific power of 440

Wh/kg and a life of I000 cycles.

ACKNOWLEDGMENT

The author thanks J. E. Harmon for editorial review of the manuscript.

work reported was funded by the U.S. Department of Energy, Conservation and

Renewable Energy, under contract No. W-31-1Og-ENG-3B.

The

312
REFERENCES

I. Tomczuk, Z., B. Tani, N. C. Otto, R. F. RoGhe, and D. R. Vissers.

J. Electrochem. Soc., Vol. 129, p. 925 (1982).

2. Kaun, T. D. J. Electrochem. Soc., Vol. 132, p. 3063 (1985).

3. Kaun, T. D., T. F. Holifield, and W. H. DeLuca. Lithium Disulfide Cells

Capable of Long Cycle Life. Ext. Abstracts, 174th ECS Mtg., p. 71,

Chicago, IL (Oct. 1988).

4. Kaun, T. D., T. F. Holifield, M. N. Nigohosian, and P. A. Nelson.

Development of Overcharge Tolerance in Li/FeS and LiFeS 2 Cells. Ext.

Abstracts, 174th ECS Meeting, Vol. 88-21, p. 67, Chicago, IL (Oct. 1988).

5. Redey, L. Proc. 22nd IECEC Mtg., Philadelphia , PA, Vol. 2, p. 1091 (1987).

6. Molyneux, J., G. Sands, S. Jackson, and I. Witherspoon. IECEC Meeting,

Philadelphia, PA, p. 975 (1987).

7. Bones, R. J., J. Coetzer, R. C. Galloway, and D. A. Teagle.

J. Electrochem. Soc., Vol. 134, p. 2379 (1987).

8. Dell R. M., and R. J. Fortes. Proc. 22nd IECEC Mtg., Philadelphia, PA,

p. 1072 (1987).

9. Tilley, A. R., and R. M. Bull. Proc. 22nd IECEC Mtg., Philadelphia, PA,

p. 1078 (1987).

I0. Sudworth, J. L., R. C. Galloway, and D. S. Dermott. Electric Vehicles,

Vol. 73, p. 14 (Autumn 1987).

II. Wright, M. L., and J. L. Sudworth. Ext. Abstr., Electrochem. Soc. Mtg.,

Honolulu, HI, p. 149 (October 1987).

12. Coetzer, J. J. of Power Sources, Vol. 18, p. 377 (1986).

13. Coetzer, J., and M. J. Nolte. U.S. Patent 4,592,969 (1986).

14. Bones, R. J., J. Coetzer, R. C. Galloway, and D. A. Teagle. Ext. Abstr.,

Electrochem. So¢. Mtg., San Diego, CA, p. 1122 (October 1986).

15. Coetzer, J., and M. M. Thackeroy. U.S. Patent 4,288,506 (1986).

16. Coetzer, J., R. C. Galloway, R. J. Bones, D. A. Teagle, and P. Mosely.

U.S. Patent 4,546,005 (1985).

The submitted manuscript has been authored

by a contractor of the U.S. Government

under contract No. W-31-| 09- ENG-38.

Accordingly, the U. S. Government retains a

nonexclusive, royalty-free licente to publish

or reproduce the published form of this

contribution, or allow Others to do so, for

I U. S. Government purposes.

313
TABLE I. COHPARISON OF HIGIt-TEMPERATURE CELL COUPLES

System

Li-A1/FeS

Li-A1/FeS 2

Li-A1/FeS 2

Na/FeCl 2

Na/NiCl 2

Na/S

Positive Average

Electrode Open

Discharge Circuit

Product Voltage_ V

Li2S,Fe 1.34

Li2FeS 2 1.67

Li2S, Fe 1.50

NaCl, Fe 2.35

NaCl, Ni 2.59

Na-S,. 5 2.01

?;

_ =: -::.

314

Theore t i cal

Specific

Energy,

Vh/kg

460

475

630

728

796

760

=

E
TABLE II. LIT_II/M/SULFIDE CELLS--PROJECTED PERFORMANCE

Characteristic

Capacity, Ah

Specific Energy, Wh/kg

Specific Power at EOD, W/kg

Weight, g

Volume, cm 3

Based on Eaton DSEP Van Requirements

Monosulfide Monosulfide Disulfide Disulfide

Multiplate Bipolar Bicell Bipolar

313 92 137 87

112 160 175 210

127 182 199 239

3410 686 1250 663

1310 250 455 240

TABLE III. LITHIUM/SULFIDE VAN BATTERIES--PROJECTED PERFORMANCE

Basis: Eaton DSEP Van

Vehicle Weight: Base, 1376 kg; load, 273 kg; total, 1649 kg + battery weight

Range: 125 miles (range at end of battery life: 100 miles)

Energy Usage: 0.194 Uh/tonne-mile (FUDS cycle)

Power Requirement (end of discharge, end of battery life): 23 kW/tonne (0-50 mph/20 s)

Electronic Control: 400 A max, 200 V max open-circult voltage

Characteristic

Monosulfide

Multiplate

Cells per Battery 150

Number of Parallel Strings 1

Capacity (FUDS Cycle), kwh 56.4

Power at EOD, i kW 64

Specific Energy, Wh/kg 83

Specific Power, W/kg 94

Weight, kg 680

C_II Fraction of Battery gt 0.75

Volume, L 415

Dimensions, in. 59Lx43WxlOH

Rate of Heat Loss, W 300

Monosulfide Disulfide Disulfide

Bipolar Bicell Bipolar

450 224 336

3 2 3

49.4 49.0 46.8

56.3 55.7 53.3

127 132 167

144 151 190

390 370 280

0.80 0.75 0.80

227 225 163

59LxlOD(3) _ 42Lx33WxlOH 42LxlOD(3) _

200 200 200

"20Z additional power over requirement is provided for new batteries to allow

for battery degradation,

bDimensions shown are for each of three modules.

315
TABLE IV. DESIGN AND PERFORMANCE DATA FOR Li-A1/FeS 2

CELL DISCHARGE FOR I-ms AT 440°C

Components:

Positive Electrode

Separator

Electrolyte

Negative Electrode

Bipolar Plate

Design Parameters:

Mass, g/cm 2

Thickness, cm

Area Loading, mAh/cm 2

Electrolyte Condition

Performance Parameters:

Measured:

Area Specific Impedance, mohm-cm 2

Area Specific Power, W/cm 2

Derived:

Specific Energy, Uh/kg

Specific Power, kUh/kg

Power Density, kU/L

Power to Energy, W/Wh

FeS 2

BN + MgO

LiF-LiCI-LiBr

Li-AI

Mo

0.292

0.123

42

flooded

41

24.1

144

82.7

196

574

316
TABLE V. ELECTROLYTE CONFIGURATION AND KEY CELL PARAMETERS

Cell Capacity @ 3-h Rate, Ah

Electrolyte

Number of Elements I I00 6

Active Length, cm 30 18 18

Outside Dimensions 3.0-cm dia 0.5-cm dia 0.4 x 9.0 cm

Wall Thickness, mm 1.5 0.4 0.4

Electrolyte Outside Area, cm 2 283 2830 1950

Current Density, 3-h Rate, mA/cm 2 118 24 34

Voltage Loss @ 1.5 C Rate at

the Following Resistivities

6 ohm-cm," V 0.50 0.03 0.04

50 ohm-cm, V - 0.23 0.31

I00 ohm-cm, V - 0.46 0.62

Positive Electrode (central sodium)

Loading Density, b Ah/cm 3 0.25 0.35 0.35

Thickness, mm i0.5 c 1.5d 2.9e

Electrolyte Configuration

Single Multiple Flat-Plate

Large Tube Tube Compartmented

I00 200 200

•Resistivity of B"-alumina at 250°C is about 6 ohm-cm

bBased on rated capacity of cell

CRadlal width of annulus

dAverage radial width of annulus (irregular shape)

"Half-thickness of flat electrode reacted on both faces

3]7
1.80

1.60

_J

> 1.40

g'l

O

:>

I .20

1.0(_

0

A

FeS2 --_

Electrolyte: LiCI-KC! Eutectic

Temperature: 417°C

Current Density: 12 mA/cm 2

I I I

(a)

I ,2c_

35

1.801 37.9 (37.5)

1.60=-_ I 42.6 (41.81

,-,I

:>

o 1.4 C

I 1 I

40 45 50

(b)

I I [

25 50 75 100

% of Capacity

B __) Li3Fe2S4 -_ Li2.xFel-xS2 + Fel-xS c

% of Capacity

55

D

Li2FeS2 --_ Li2S + Fe

FIGURE 1. DISCHARGE CURVE FOR LiAi/FeS 2 CELL.

318
o,o 0

0

(A) eBe_lOA Ile:_

u_

t_

Z

m

mO

w...I

>-,_

Oe_

0 _

u.Z

>.m

I-.w

--a

unl

w_

__1

O--

w

m

u_

319
e, :

o

..o

:o

it) iII:

m

I•

ID

LN

• U. • &. • m " / : / - / • /

io"

/

•I/

_®,.. •

_

i ,,,! o o o o

('deo "oaql

A0 -.

m0

"_o

0--

_(I) ,_

-,%;>,,-

.__!1o

o_=___-

_-- _),'T;, t--

P_'5lo

%) u0!;eZ!l!ll'l

ooooo

w

L_i

w_

__j

Z_

_L

_, zo

o

320

=E
l .._ 0I

.,..1

.,..I

ILl

L.."

• ...•

. L."

• lb. '

r'_.

n,'.

", L.I

_?• . lb..

'. Lo

•.o

0

W

0-

o):_

ILl

p_

t_0

#

ttl UZ0U

..J

..J

UJUW

r_m

..J Z0

n"mm

-r

F"m

..J

..J 0m

LL

321
C0

0

0

>

Z

=.

4,

p-i

÷

I--

322

Zi

_m

,=J W

_U

"I-,IU>. U'} --0

r_

ULJ I

ml

_Z

OI-

_/.iW

i-u.i

i

LL IE
r

0

r')

• OOCX_

_:_0 0 0 00q.

.... i .... 1 ....

0 0 0

C%i T0

c_J

_ooo_ _,

.... I .... I .... I ....

Le) 0 _ 0

st.. ,r" 0

(A) e6e;IOA lie:)

co

ulZ ul mlUm inlm

m

P

m

.<_

_j

- _ ___z

" _ {2.

- o.<_, <

<'_u"

-s" _"s __goO'S

0

m

Lr) u- _"

zull-

,,,_<

I-- a. E

u

t.

m

323
FIGURE 7. PROPOSED FLAT-PLATE

/_'-ALUMINA ELECTROLYTE FOR

SODIUM METAL CHLORIDE CELL.

324

=

=

=
N90-20481

ADVANCES IN AMBIENT TEMPERATURE SECONDARY LITHIUM CELLS

S. Subbarao, D.H. Shen, F. Deligiannis,

C-K. Huang and G. Halpert

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, California 91109

JPL is carrying out a NASA/OAST sponsored R & D program on the

development of ambient temperature secondary lithium cells for

future space applications. The goal of the program is to develop

secondary lithium cells with a i00 Wh/kg specific energy and

capable of i000 cycles at 50% DOD. The approach towards meeting

these goals initially focused on several basic issues related to

the cell chemistry, selection of cathode materials and electrolytes

and component development. We have examined the performance

potential of Li-TiS2, Li-MoS3, Li-V6013 and Li-NbSe 3 electrochemical

systems. Among these four, the Li-TiS_ system was found to be the

most promising system in terms of achlevable specific energy and

cycle life. Major advancements to date in the development of LiTiS 2 cells are in the areas of cathode processing technology, mixed

solvent electrolytes, and cell assembly. This paper summarizes

these advances made at JPL on the development of secondary lithium

cells.

INTRODUCTION

Ambient temperature secondary lithium batteries have several

intrinsic and potential advantages including higher energy density,

longer active shelf life, and lower self discharge over

conventional Ni-Cd, Ni-H2, Pb-acid and Ag-Zn batteries. Successful

development of these batteries will yield large pay-offs such as

2-3 fold increase in energy storage capability and a longer active

shelf life of 2 to 4 years over these systems. A detailed analysis

of the strengths and weakness of secondary lithium batteries has

indicated that these batteries are suitable for applications

requiring less than 1 kW power and limited cycle life. Some of the

projected space applications of these batteries are for Mars Rover,

planetary space craft/probes, astronaut equipment, and

geosynchronous spacecraft.

Several ambient temperature secondary lithium systems such as:

Li-TiS2, Li-MoS2, Li-NbSe3, Li-Mn02, Li-V205 etc. are currently under

development for consumer and defense applications (i). Most of

these systems are still in the research stage while a few of them

are commercially available in small capacities (less than 1 Ah).

Performance of the cells that are commercially available fall short

325
of meeting the space application requirements (2).

Under a NASA OAST sponsored program, Jet Propulsion Laboratory is

developing ambient temperature secondary lithium cells for future

space applications. The primary objective of the program is to

develop ambient temperature secondary lithium cells with I00 Wh/kg

specific energy while delivering I000 cycles at moderate depths of

discharge (50%). The major thrusts of this program are to improve

cycle life and safety of these cells. Approach has involved work

in three areas: i) chemistry and material research, 2) component

development and interactions, and 3) performance and safety. This

paper describes the advances made at JPL to date in these areas.

CHEMISTRY AND MATERIAL RESEARCH

Attempts to improve the performance capability of the state of art

secondary lithium cells has been inhibited by the unavaiiabiilty

of materials with the desired properties and poor understanding of

the cell chemistry(3). For these reasons the major emphasis of

this effort has been directed towards identifying candidate

electrode materials and a stable electrolyte and developing

fundamental understanding of materials and cell chemistry. Some

of the specific objectives in this area are : i) Selection of a

cathode material with a theoretical specific energy greater than

400 Wh/kg providing more than I000 cycles, 2) Identification of an

electrolyte that is stable towards lithium and has a conductivity

greater than 10 .3 ohm-l-cm -I, 3) Evaluation of the use of lithium

alloys and other materials with low equivalent weight (40 gm/eq)

to extend the cycle life performance andsafety of the cells and

4) Development of concepts for overcharge/overdischarge protection

of cells. Significant progress was made by JPL in the first three

above topics as described below.

Cathode Materials

Cathode materials required for these cells must possess several

desirable properties such as high specific energy, good

electrochemical rechargeability, good electronic conductivity, high

lithium diffusivity, and chemical compatibility towards

electrolyte. Several materials have been reported in the

literature for use in secondary lithium batteries (3). These

materials can be classified into two categories: transition metal

chalcogenides (TiS2, MoS3, MoS2, NbSe , etc ) and transition metal

oxides (VzO5, V6013 , MnO2, CoO2,etc .). Lithium 3 " is inserted into these

host structures electrochemically during discharge (Figure I).

Ideally, the host lattice may undergo Only a minor structural

change during lithium intercalation/insertion process (discharge),

and return to its original state after charging process (the

deintercalation). Four of these materials ( TiS 2 MoS3, V6013 , and

NbSe 3 ) were selected for detailed examination. These four

materials were evaluated for their performance (specific energy,

and cycle life) in experimental cells. In order to overcome poor

326
cycle life performance of the lithium electrode, a high lithium to

positive electrode ("cathode") ratio (6:1) was used in these

experimental cells. Major emphasis of this study was focussed on

determining the cycle life capabilities of these cathode materials.

The results of these studies are summarized in Table I. The

results clearly indicate that the most promising cathode material

is TiS 2 in view of its long cycle life performance capability and

realizable specific energy. The second material of choice is

NbSe 3. Even though MoS 3 and V6013 have higher theoretical specific

energy, the materials showed poor cycle life performance and low

useful specific energy. These studies led to the selection of Tis 2

as the candidate cathode material for further studies with

electrolytes and anode materials.

Electrolytes

The ideal electrolyte required for ambient temperature secondary

lithium cells must exhibit high conductivity, wide electrochemical

operating window, high stability towards lithium, and compatibility

with cathode materials and separator, low density, and low

viscosity. Some of the electrolytes that have been investigated

for this application are THF, 2-MeTHF, Methyl Formate, and

Propylene Carbonate (4). These single solvent electrolytes provide

only limited cycle life capability because of their reactivity

towards lithium. It has been suggested (5) that the use of mixed

solvent electrolytes may alleviate this problem due to the

formation of a beneficial passivating film on the lithium anode and

the synergistic effects associated with the use of the selected

mixed solvents. A number of mixed solvent electrolytes are

currently being investigated in our laboratory for use in Li-TiS 2

cells. These electrolytes contain 2-MeTHF and THF as the base

solvents and Ethylene Carbonate(EC), Propylene Carbonate (PC) and

3-Methylsulfolane(3-MeS) as the co-solvents. The important

properties such as conductivity, stability towards lithium and

lithium cycling efficiency of these mixed solvent electrolytes were

evaluated . Conductivity of the electrolytes was determined by

standard methods. Stability of the electrolytes towards lithium was

investigated by microcalorimetry. The heat evolved from the

uncycled experimental cells in open circuit condition was measured

at 25°C with a heat conduction Hart Scientific microcalorimeter.

The cycling efficiency of lithium in various electrolytes was

determined by cycling experimental Li-TiS 2 cells. The lithium

cycling efficiency/ figure of merit (FOM) was calculated using the

relationship:

F.O.M.L_

total accumulated discharae capacity

theoretical Li capacity

Comparative performance of selected electrolytes is summarized in

Table II. Among the co-solvents investigated, EC appears to be a

327
promising co-solvent. The mixed solvent electrolytes containing

EC (EC/THF, EC/2-MeTHF, EC/THF/2-MeTHF, EC/3-MeS/2-MeTHF) showed

higher conductivity and improved stability towards lithium compared

to the electrolytes containing no EC (based on open circuit stand

tests and microcalorimetry). The AC impedance and FTIR studies (6)

indicated that the improved performance of EC/2-MeTHF electrolyte

may be due to the formation of a thin LiCO_ film on the surface of

the lithium electrode. From the results it can also be observed

that EC containing electrolytes yielded higher FOM values than

other electrolytes. A FOM of 38.5 was obtained with 10%EC+90%2-

MeTHF electrolyte and this is 33% higher than the FOM of the bench

mark 2-MeTHF electrolyte (29.5). Although the achievement of this

FOM represents a significant advancement, the FOM value presently

remains lower than the target objectives ( FOM ~ 80) . Work is

continuing to identify an electrolyte that can meet program

targets.

Lithium Alloys

The limited cycle life performance of ambient temperature lithium

cells is primarily attributed to the reactivity of the lithium

anode with organic electrolytes. One method for reducing the

reactivity of the lithium anode towards the electrolyte is by the

use of lithium alloys as anode materials. Further, the use of

lithium alloy anodes may also improve the safety of the cells. Six

lithium alloys were selected for preliminary experimental

evaluation after a detailed review of the literature. These alloys

are Li-A1, Li-Si, Li-Sn, Li-Zn, Li-Pb, Li-Cd. All these materials

are basically lithium rich alloys. These materials were selected

based on their electrochemical potential, equivalent weight,

lithium diffusivity, reversibility and thermodynamic stability

towards organic electrolytes. Some of the important properties of

these alloys are summarized in Table III. Experimental evaluation

of these alloys was initiated on these anodes and is in progress.

COMPONENT DEVELOPMENT AND INTERACTIONS

Component characteristics and cell design also play a significant

role on cell performance. For this reason, work is in progress on

the fabrication of TiS_ cathodes and design analyses. Details of

the progress made in these two areas are described below.

TiS 2 Cathode Processing

The rate capability and faradaic utilization of TiS 2 cathode are

dependent on the chemical composition of TiSz, nature of the

binder, additives/ conducting diluents, as well as thickness and

porosity of the electrode. TiS 2 processed by the vapor transfer

method was found to show higher faradaic utilization compared to

the TiS z materials procured from commercial sources (Table IV).

This improved performance was found to be associated with purity

328
and stoichiometry of the material. Hence, it was decided to use

TiS2 prepared in house by the improved vapor transport method in

our future studies.

During cycling the cathode undergoes a volume change due to the

intercalation/deintercalation of Li into/from TiS 2 crystal lattice.

The electrode structure must be capable of accommodating these

changes, otherwise it will disintegrate during cycling. Studies

carried on various binder materials identified EPDM as the most

suitable binder material for the construction of Tis2 cathodes (7).

Three methods : brushing, rolling, and presslng methods were

examined for the fabrication of TiS2 cathodes. The brushing method

was found to be suitable for the fabrication of small electrodes

in the glove box. The rolling and pressing methods offered several

advantages including amenability to scale up, and uniform loading.

Typical characteristics of the electrodes fabricated by various

methods is given Table V. Work is in progress to understand the

relationships between electrode properties and performance.

Design Trade off Studies

Some of the important parameters to be considered in the design of

secondary lithium cell are ratio of electrode capacities, operating

current density, electrolyte composition and quantity, case and

grid materials and cell size. A computer program was developed (8)

to understand the influence of these parameters on specific energy

and to optimize the cell design. The results of these studies

indicated that a specific energy of 80-100 Wh/kg is achievable for

high capacity cells( >20 Ah ). The negative to positive electrode

capacity ratio has minimal effect on specific energy of the cells.

A high anode to cathode capacity is required to account for the

degradation of Li electrode and to achieve maximum cycle life.

Cells cannot operate at current densities higher than 2 mA/cm 2

because of the poor conductivity of the electrolyte. The cell case

is the major contributor to cell weight( Fig. 2 ). Among the

active materials, Li contributes the least to the overall weight.

The collector grids contribute more to the weight than does lithium

alone. The use of titanium cases and aluminum grids will reduce

cell weight considerably. Experimental work is in progress to

determine the influence of some of the important parameters on the

cycle life performance of the cells.

CELL PERFORMANCE AND SAFETY

An assessment was made of the status of secondary lithium cells by

experimental evaluation of Li-MoS2, Li-NbSe 3 and Li-TiS 2 cells. LiMoS. and Li-NbSe 3 cells (AA) were procured from the industry. The

Li-TiS 2 cells were fabricated in house with non-optimize and cell

hardware. These cells were evaluated for their charge/discharge

characteristics, rate capability and cycle life. Performance

characteristics of these cells is summarized in Table VI. These

32g
cells have a cycle life of 100-250 cycles at moderate discharge

rates ( C/5 ). Li-NbSe 3 cells exhibited high specific energy and

Li-MoS 2 cells showed long cycle life. The low specific energy and

cycle life of 5 Ah cells may be due to the non-optimize cell design

and cell hardware. Some of the important problems identified are

formation of soft shorts during cycling at low discharge rates (

C/lO ) and inferior performance at low depth of discharge in LiMoS2 system. Evaluation of the safety of these cells is in

progress.

CONCLUSIONS

Li-TiS 2 system appears to be the most promising system, among the

ambient temperature secondary lithium sys£ems examined, in view of

its higher realizable specific energy and cycle life. Use of mixed

solvent eiectrolytes was found to Improve the Cycle life of this

system. Among the various electrolytes examined, 1.5 M LiAsFj EC

+ 2-MeTHF mixed solvent electrolyte was found to be more stable

towards lithium. Experimental cells activated with this

electrolyte exhibited more than 300 cycles at 100% DOD. Design

trade off studies of this system indicated that a practical

specific energy of 80-100 Wh/kg is achievable in 20 Ah cells and

higher. The technology of these ceiis is still far from meeting

the requirements of planetary space missions. Significant progress

is still needed, particularly with respect to cycle life, tolerance

to over-charge and over-discharge, safety and cell size. Many of

these key problems require significant scientific and technological

innovations in the areas of electrolyte technology, alternative

lithium based anode materials, new additives and cell design.

ACKNOWLEDGEMENTS

This Work represents one phase of research performed by Jet

Propulsion Laboratory, California Institute of Technology sponsored

by the National Aeronautics and Space Administration under contract

NAS-7-918.

330

=

z
REFERENCES

1) B. B. Owens, "Survey of Lithium Ambient Temperature Secondary

Systems", Fifth International Seminar on Lithium Battery

Technology and Applications, Deerfield Beach, Fla, March 1989.

2) S. Subbarao et al, "Ambient Temperature Secondary Lithium

Batteries for Space Applications", Proceedings of the Fourth

Annual Battery Conference on Applications and Advances, Long

Beach, CA, January, 1989.

3) K. M. Abraham, "Status of Rechargeable Positive Electrodes

for Ambient Temperature Rechargeable Batteries", J. Power

Sources 7, 1(1981).

4) V. R. Koch " Advances in Organic Electrolytes for Rechargeable

Lithium batteries", Fourth International Seminar on Lithium

Battery Technology and Applications, Deerfield Beach, Fla,

March 1988.

5) S. Tobishima, M. Arakawa, T. Hirai and J. Yamaki, Journal of

Power Sources, 20, P. 293-297 (1987).

6) D. H. Shen, S. Subbarao, F. Deliagiannis, and G. Halpert, "

Evaluation of Mixed Solvent Electrolytes for Ambient

Temperature Secondary Lithium Cells", Proceedings of the 33rd

International Power Sources Symposium, p.45, (1988).

7) S. P. S. Yen, D. H. Shen, and R. B. Somoano, Journal of

Electrochemical Society, 130, P.5, May (1983).

8) D. H. Shen, S. Subbarao, S. P. S. Yen and R. B. Somoano,

Journal of Power Sources, 18, P127-131 (1986).

331
11.

0 w

z_"

111

"7O

0

111

O0

>-0

0 0

0 tO I_ 0

0 I I 0

¢ID 0 0 _-

A '_ _ A

I_ W wa.

_r_ zx

UJ UJ

C) O

0 _

b..

F-- _d

Iii

0

Z

'< a

n-bOo <

ILl i-

> _i ,-: oi

0

i

UJ

_J

rn

I--

A

Z

It)

i

I-- _ O0 0 C/_

=_ ic- =_ >® z

332
00

LLI

1---

>- O

F--

Z

>- >-

A

z

_0

.j ii1

owo

LL.I __ i_._ 0

LIJ 0

_0Waiaiiii

l--

rr

t.l..I

O_ 0

rr

O.

i

l.lJ

....I m

<_

t--

N

ko0 0

Ill

k-

>-

_J0

r_

ILl

--I

uJ

0

LL.

7"

k333

cO

(D

_D ,4,

O.I

O

O

LI.,

7"

F--- t--io

OOO 00

:z::+++_w

O r,D (DI :Z:::

LU tJ.! O. c'_ FI O O O _ O

CO TM

00

n-'

I--

LLI II

o_0

_

O_z

rr"

OOO

LLI

Z

oa:

O

ILl
>.

O

<

i

(n

m _m_

<

_J

m f

m

>. ILl o

-J

r_ <

<_

F4K 41 _ 4_ G

O O O O 'O

_- ,- c: C; O

U_

>.

.J

r_

m

.J

334
n"

LLI

I-- <<

rt"

O9

ILl m

InLO

O. 0

rr

O.

LLJ

_J

rn

I--

Z

A

_ood

.<l--ILl

rr_

>.

I-ra

09 --_ Z o

I.Uo

r,. <

LL] m

I--

rrni

>..

rr

l---

I.IJ 0I0 09 0 _"

hO9

LU

(J

rr0

_O

O_0

_000

O_

..J

-3

_t6m0

r_.

o_

q-. <

09

09

(5

LU

C3

0

_D

cO0

O90

cO

C)<

rr

UJ0

335
CO

LUa0

F-<0

LL0

O9

LUn

Frr

LUn0

nO.W

_.J

rn<

F-

(97

09

09 cy in

ILl -I- 4-

n" tO ,_.

(9Z

_J 4-+

0o _o

rr" o_ to ZmI

O9 cy

rr_ co

rn ,- ,_-

++

_)

cY+

+

eo+÷

E

•-0

>. 09 _-

09 >" >-

ocw co_

wz _z_

o. _"a0<

00< rr O.

rr "_00<

O. I- ._ O.0

336
ii_

0

Z

0

i

,_i

0

ill

rligRime

I.L,

..t

,,j

Dm

C

C

C

)

3

J

Dm

iilii

i.I

C

i

o_

I--

!

()

J

f

0

0 •

)

iim liB

u) l-i _.1

()

()

337
W

a

m

ILl

..J

_J

W

0

I

w_

._J

N

b_

I I I

o 0 o o

'q- co c',J _-

f

<

0

o

rn

<

Q

rr

o

a:

h-

,o

d

LU

0

Z

<

LU

C_

o

I

F-

<

<9

FZ

W

Z

0

O_

Z

0

0

_J

__J

W

©

_l_N3Okl3d _.LH©I3M

338
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13. 1969

ADVANCED CONCEPTS

CHAIRPERSON: A.J. APPLEBY

CO-CHAIRPERSON: MICHELLE MANZO

APRIL 12,198g

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13. 1969

SUMMARY

ADVANCED CONCEPTS

• COLD FUSION

LAST YEAR'S ADVANCED CONCEPTS

• HIGH TEMPERATURE BATTERIES:

-COMPOSITE SOLID ELECTROLYTES

• BIPOLARLi AI- Fe$2, Li AI- S, (?) Li- 02

- NEW MATERIALS

• HIGH TEMPERATURE FUEL CELLS

- MONOLITHICAND PSEUDOMONOLITHIC

- NEW ELECTROLYTES

• LOW TEMPERATURE SYSTEMS

- ULTRACAPACITORS

- NEW Li SYSTEMS

• NEW COUPLESANDDESIGNS SFRT89 O013

339
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

APPLICATIONS WILL DETERMINE WHICH

ADVANCED SYSTEM IS BEST

• FOR LEO, REQUIREMENTS ARE

- HIGHENERGYDENSITY

- HIGHCYCLE LIFE (>30,000CYCLES)

- HIGHPOWERDENSITY

POfiS/BLE(_/_ND|DATES: SOLIDELECTROLYTE SYSTEMS; RFCI

• FOR PLANETARY ROVER, REQUIREMENTS ARE

• HIGHENERGYDENSITY

• VERY HIGH POWERDENSITY

• 1000 CYCLES

POSSIBLE CANDIDATES: :LIAI-FeS2, FUEL CELL, Na/BETA/METAL CHLORIDES

• FOR SPACE PROBES, MANPACK, SMALL GEO SPACECRAFTS

. HIGHENERGYDENSITY

- SMALLSIZE

• 1000 CYCLES

- LOW TEMPERATUREOPERATION (5-20"C)

POSSIBLE CANDIDATES: AMBIENT TEMP. LITHIUM; POLYMER ELECTROLYTE SYSTEMS

• FOR MOON BASE OR MARS OUTPOST, REQUIREMENTS ARE

- EXTREMELYHIGH ENERGYDENSITY

- EXTREMELYHIGH RELIABILITY

POSSIBLE SYSTEMS: NUCLEAR; RFC 5ERT$9_OtJ

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

COLD FUSION

IF TRUE, WILL RADICALLY CHANGE THINKING

BASELOAD FOR LUNAR BASE

POSSIBLY PROPULSION: SPECIFIC POWER IS KEY; ESPECIALLY IF NO SHIELDING REQUIRED

MUST GIVE HIGH-OUALITY HEAT (300'_'CFOR RANKINE, PREFERABLY 1000 +'C, BRAYTON, OTHERS)

IMPLEMENTATION OF 1000'C, MOLTEN 6LID?

SMALL SYSTEMS; STIRLING, Na HEAT ENGINE, POSSIBLY THERMIONICS, THERMOELECTRICS

CLASSICAL CYCLES SINCE NOY- OR n - TO ELECTRICITY DEVICES EXIST

QUESTIONS: STANDBY POWER CAPABILITY?

LOAD-LEVELING BATrERIES OR REGENERATIVE FUEL CELLS WILL BE

NECESSARY, OR OFF-PEAK POWER CAN BE USED FOR PROPELLANT

MANUFACTURE ,

• WHILE EXCITING, IF TRUE, WILL TAKE A LONG TIME TO IMPLEMENT. SO WE NEED TODAY'S

SYSTEMS AND LAST YEAR'S NEW COI_CEPT_. R[SK LEVEL - VERY HIGH: REQUIRES A CONRRMATION

IF APPUCABLE, SYSTEMS & MATERIALS DEVELOPMENT

TIME FRAME: 2015. +

EFFO_I': AS RSSION

340
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

HIGH TEMPERATURE BATTERIES

• COMPOSITE SOLID ELECTROLYTES

- IMPROVEMENTSIN NaS ANDOTHERSYSTEMS(Na/TRANSITION

METALCHLORIDES,Li SYSTEMS)

- REQUIRENEW GEOMETRIES,BETTERPACKAGING

- LOWERTEMPERATURESOF CHLORIDESYSTEMSMAYGIVE

BETTERLIFE

- SHOULDBEVERYSAFE(INSOLUBLECHLORIDESIN NaAICI4;

NaFeCI4UNSAFE)

• LIAI-FeS2 BIPOLAR: 210 Wh/kg

REQUIRE Mo-PLATEDBIPOLARCOLLECTOR,SEALS

' Li/02 (?)

• LIAI-S, SOLID ELECTROLYTE: 300 Wh/kg

• NEW HIGH.TEMPERATURE MATERIALS, e.g., COATED CARBONCARBON (INSULATING OR CONDUCTING, HIGHLY CORROSION

RESISTANT_ SErRTS§-O015

.SpACE, ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13. 1989

HIGH TEMPERATURE FUEL CELLS

. MONOLITHIC AND PSEUDOMONOLITHIC STRUCTURES

(EXTRUSION, ETC)

, MONOLITHIC ELECTROLYZERS (MAY USE WITH LOWTEMPERATURE FUEL CELL)

• LOW-TEMPERATURE SOLID ELECTROLYTES FOR FUEL CELLS

SERTag-o01 6

341
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1888

LOW-TEMPERATURE SYSTEMS

• 20,000 J, 5 Wh/kg ULTRACAPACITORS

- HIGHERVOLTAGE REQUIRED(ORGANICS,LOW TEMPERATURE

MOLTEN SALTS)

" HIGHERFARADAICPSEUDOCAPACITANCE(250 iJF/cm2)

• HIGHERAREAS

• LI POL-YMER ELECTROLYTES

- NEWSTRUCTURES(OTHERTHANPEO/CF3SO3-)

• AUTONOMOUS Li CONDUCTORS

• LI - METAL OXIDE CATHODES (C002, MnO2)

- PROBLEMOF CHARGE POTENTIALWINDOW

• NEW AQUEOUS COUPLES (H2-BilMnO2 , etc)

• NEW DESIGNS WITH OLD COUPLES (BIPOLAR, e.g., WITH Ni/H2)

• RELIABLE, TRUE "SINGLE-CAN" REGENERATIVE FC/ELECTROLYZER

_ SPACE ELEcTROCHEM_AiL. RESEARCH& TECHNOLOGY CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13. 11_8B

GENERAL COMMENTS

• RELIABILITY vs WEIGHT & COST TRADE-O_FFFOR SPACECRAFT

HIGHCAPABILITy(LOWWEIGHT,VOLUME,ETC.) MUST BE CONSISTENT

WITH RELIABIL!TY

• MOST FUNDS ARE SPENT ON EXlSITING SYSTEMS. NEW SYSTEMS

NEED A RELIABILITY ANALYSIS BEFORE FUNDS ARE COMMII-rED FOR

PILOT PRODUCTION OF INAPPROPRIATE DESIGNS FOR MISSION

• ACCELERATED TESTING REQUIRED: WEIBULL AND OTHER ANALYSIS

• RISK LEVEL

• IS 1HEREA WAYOUT iF THEAPPROACHDOES NOTWORK? ARETHERE

ALTERNATIVEOPTIONS?

" OPTIONbAVOIDRISK

• TIME FHAME, LEVEL OF EFFORT

- VARIABLEFROMTRANSFERTO PRIVATESECTOR;SMALL EFFORTTO

(FORFUSION,IF TRUE) MASSIVEPROGRAM

342

SERT89 6Ot IB

z
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

HYDROGEN-OXYGENFUEL CELLS AND ELECTROLYZERS

CHAIRPERSON: JAMESHUFF

CO-CHAIRPERSON: PAULPROKOPIUS

APRIL12, 1989

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1guy

HYDROGEN-0XYGEN FUEL CELLS AND ELECTROLYZERS

FUTURE MISSIONS BENEFITED

• NEAR-EARTH

• PATHFINDER

• UNSPECIFIED - AS YET UNDEFINED

343
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1089

NEAR-EARTH MISSIONS

• INTEGRATED FUEL CELL - ELECTROLYZER

_g 5FRT D2

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

NASA LEWIs RESEARCH CENTER APRIL 11-13. 1989

PATHFINDER

LONG-LIFE, RELIABLE, EFFICIENT PEM AND ALKALINE FUEL CELLS

AND ELECTROLYZERS

344

JHRg-SERTo]
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

UNDEFINED MISSIONS

• LONG-LIFE, RELIABLE, EFFICIENT, COMPACT PRIMARY OR REGENERATIVE

FUEL CELL SYSTEMS

_8_ SERT_4

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, tg8g

THESE SYSTEMS ARE ENABLING FOR THE INDICATED MISSIONS

JH_J-SERTOS

345
SPACEELECTROCHEMICAL RESEARCH& TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1_8g

TECHNOLOGY ISSUES

• COMPONENT LIFE TESTS IN FULL PLANAR AREA HARDWARE

-- FAILURE MECHANISMS

-- REACTANT CROSSOVER

-- CORROSION

--DEVELOPMENT OF PREDICTIVE MODELS

,TH89 SERT 06

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, lgSg

DEVELOPMENTAL NEED

• PASSIVE SYSTEM OPERATION FOR HIGH RELIABILITY

_rlSg.sERT 07

346
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL | 1-t3, 1988

RESEARCH ISSUES

NO PROVISION FOR LONG-TERM RESEARCH IN MISSION DRIVEN PROGRAMS

._Jg.SERT o_

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

NICKEL ELECTRODE WORKSHOP

CHAIRPERSON: ALBERT ZIMMERMAN

CO-CHAIRPERSON: MARGARET REID

APRIL 12, 198g

MR_ SER101

347
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, lgBg

ADVANCED NICKEL ELECTRODE TECHNOLOGY

• LIGHTWEIGHT NICKEL ELECTRODES (OPTIMIZED FOR HIGH ENERGY

DENSITY)

• SINTERED

• FIBER

• NICKEL ELECTRODES OPTIMIZED FOR LONG CYCLE LIFE

• OPTIMIZED SUPPORTINGPOROUSSTRUCTURES

• TECHNOLOGY MAINTENANCE

MR_S[RT 0"2

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY CONFERENCE

NASA LEWIS RESEAR_ENTER APRIL 11-i_, |989

APPLICATIONS

• LIGHTWEIGHT Ni ELECTRODES

• MODERATEcYCLE LIFE (_ 1000 CYCLES,> 50% DOD)

• HIGHENERGYDENSITY(~ 100 Wh/kg)

e MARS ROVER

• LONG-LIFE N1 ELECTRODES

• > 30,000CYCLES (50% DOD)

• SPACE STATION

• 10-20YR (GLOBAL CH'ANGETECHNOLOGYMISSIONS)

• TECHNOLOGY MAINTENANCE

• ASSUREHIGH RELIABILITYFOREVOLVINGTECHNOLOGY

348
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1999

CRITICAL TECHNOLOGY NEEDS

• STRUCTURAL INSTABILITY

• SWELLING

• EXTRUSION

• LOSS OF OHMIC CONTACT WITH CURRENTCOLLECTOR

• IMPROVED ACTIVE MATERIAL MECHANICAL PROPERTIES

• ADDITIVES

• ELECTROLYTE

• IMPROVED TEST PROCEDURES (TECHNOLOGY SUPPORT)

• ACCELERATED/REALTIME

• LOT CONFIRMATIONTESTING

• ROUTINEDPA OF CELLON LOT BASIS

• MODELING

• INPUTS- IMPEDANCE,PORE DISTRIBUTIONS,ETC.

• ANALYTICALMETHODS

e CORRELATION OF ELECTROCHEMICAL AND STRUCTURAL CHARACTERISTICS btR89.SERT04

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13o 1989

RESEARCH NEEDS

• CAUSES FOR STRUCTURAL INSTABILITY

• ACTIVE MATERIALMOVEMENT/SHEDDING

• SINTERFRACTURING

• SHORTING: DETECTIONAND CAUSES

• CORROSION RATES

• INHIBITION

• ACCELERATINGFACTORS

• ADDITIVES

• IMPROVEDCHARGEEFFICIENCY

• IMPROVEDCAPACITY(INCREASEDUTILIZATIONOF Ni)

• ELECTROLYTE CONCENTRATION - EFFECTS ON PERFORMANCE

• INTERACTIONS WITH HYDROGEN GAS

• STORAGE

• SELF-DISCHARGE

• EFFECTON CAPACITY,SECOND PLATEAUEFFECTS .,. s_.'ro5

349
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1080

RESEARCH NEEDS

TESTING

• CHARGE CONTROL METHODS

SCALE-UP

• EVALUATEPRESENTASSUMPTIONSOF LINEARSCALING

• EXAMINELIFETEST DATAAND FAILUREMODES

• FUNDAMENTAL ELECTROCHEMISTRY OF NI ELECTRODE PROCESSES

• APPLICATIONOF NOVELANALYTICALTECHNIQUES

(CALORIMETRY,NEUTRON SCATI'ERING, ACOUSTIC EMISSION,

NOISESPECTRUMANALYSIS,EXAFS)

MRg'_SE/_TO_

SPACE ELECTROCHEMICAl. RESEARCH & TECHNOLOGY

.... CONFERENCE _ _ _ _

NASA LEWIS RESEARCR (3ENTER APRIL 11-13, 1989

E C:c

SUMMARY

• CELLS CONTAINING Ni ELECTRODES ARE ONLY ESTABLISHED

LONG CYCLE LIFE TECHNOLOGY

• RELIABLE

• SAFE

• _MPROVEMENTS SHOULDBESUPPORTED ANDAREACHIEVABLE

• OPTiMiZATiON FORH_GHENERa_' DENSm'

• OPTIMIZATIONFOR LONG CYCLELIFE

• TECHNOLOGY MAINTENANCE PROGRAM _'HOLILD BE INSTITUTED

• ASSURECONTINUEDRELIABILITY

• SMOOTHLYTRANSITIONADVANCEDTECHNOLOGYINTO PRACTICE

MA_ $_AT 07

3_
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 1 !-13, 1989

WORKSHOP ON

ADVANCED RECHARGEABLE

BATTERIES

G. HALPERT AND Jo SMITHRICK

SPACE ELECTROCHEMICAL R&T

CONFERENCE

APRIL 12, 1989

GWJ580 SERT 01

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1080

PRIORITY

OUTER

PLANETARY

ORBIT

APPLICATIONS

CHARGE/ TYPICAL TYPICAL

DISCHARGE OPERATIONAL PEAK POWER AND

DURATIONS CYCLES REQUIRED ENERGY STORAGE REQUIRED

ACTUAL QUAL' DESIRED (5 MINUTES)

C- 2 hr 500 1,000 2,000 0.5 C (I kWh)

D - 0.7 hr

INNER C - 2 hr 3,000 6,000 10,000 1.8 C (2 kWh)

PLANETARY D - 0.7 hr

ORBIT

GEO C - 22.8 hr 1,500 2,000 4,000 1.S C (S kWh)

D - 1.2 hr

PLANETARY C - 12 hr 300 600 800 1.3 C (3 kWh)

ROVER D - 3 hr

LUNAR C - 11 DAYS 80 160 350 0,02 C (5 MWh)

BASE D - 17 DAYS

LEO C- 1 hr 30,000 35,000

D - 0.6 hr

GEO = GEOSYNCHRONOUS ORBIT

LEO = LOW EARTH ORBIT

'OUAL = MINIMUM NUMBER OF CYCLES NEEDED TO OUALIFY FOR APPLICATION

50,000 1.1 C (25 kwh)

C,_,F_FSE RT 12

351
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

OTHER NASAAPPLICATIONS

• FREE FLYERS

• ROBOTICS

• OMV

• ASTRONAUT EOUIPMENT

• PENETRATORS/PROBES

• TETHERED S/C

G_JSS_ 5ERT O_

SPACE ELECTROCH-_MiCAL RESEARCH & TECHNOLOGY CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

ASSUMPTIONS

• NO COST CONSTRAINTS

• TECHNOLOGY DEMONSTRATION - LEVEL 5 ASSUMES LIMITED

TESTING TO VERIFY CAPABILITY

• BIPOLAR TECHNOLOGY - NOT A LIMITATION

• SPECIFIC ENERGIES BASED ON BATTERY CAPACITY

352

GHJJSe9SEAT 07
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

CANDIDATES FOR LARGE POWER SYSTEMS (>2.5 kWh)

"BEST CASE ANALYSIS"

PROJECTED YEAR FOR PROJECTED

SPECIFIC ENERGY TECHNOLOGY

SYSTEM (Wh/kg) APPLICATIONS DEMO (LEVEL 5)

Na - S 150 4,5 2000

Na - S 150 3,6 2010

Ha • S 200 4,5 2010

Na - S 200 3,6 2015

u (A_).FeS2 175 4,5 2000

LI(AI)" FeS2 175 3 2010

Na - Me/x 160 4,5 2000

Ha - MCIx 160 3,6 2010

AdvN]-H2 70 3,4,5 1995

AdvNI-H2 70 6 2000

GHtJS6g SERT 09

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1989

CANDIDATES FOR SMALL POWER SYSTEMS (<2.5 kWh)

"BEST CASE ANALYSIS"

PROJECTED YEAR FOR PROJECTED

SPECIFIC ENERGY TECHNOLOGY

SYSTEM (Wh/kg) APPLICATIONS DEMO (LEVEL 5)

LI - SOLID CATHODE 100

m%,Nb%,%%)

(SAME AS ABOVE)

U - SOLID CATHODE

(CoO2)

I 1990

100 3 1995

200 1,3 2005

LI - P.E.- SOUD CATHODE 200

(SAME AS ABOVE) 200

'P.E. = POLYMERIC ELECTROLYTE

1,3 2005

2 2010

GHI.f_9 SERT 0_

353
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13o 1080

TECHNOLOGY ISSUES

Na-S (200 Wh/kg)

• REDESIGN CELL/BATTERY (200 Wh/kg)

• CORROSION FOR LONG-LIFE APPLICATIONS (>2 YRS)

• ELECTROLYTE SEALING

• CONTROL OF FAILURE MODE

• FREEZE/THAW

• LARGE SYSTEMS (2.5-100 kWh)

• STATE-OF-THE-ART - 140 Wh/kg

• PROBABILITY OF SUCCESS 80%

GH/J¢_89SERTtO

SPACE ELEC_'RQCHEMICAL RESEARCH-& TECHNOLOGY

CONFERENCE

NASA LEWIS R_ESEARCH CENTER APRIL 11-13, 1989

" TECHNOLOGY _IS_ES

U.P. LI-(AI)-FeS 2 (175 Wh/kg)

• CELL REDESIGN

• CORROSION FOR LONG-TERM APPLICATION

• HERMETIC SEAL FOR CELL

• ELECTROLYTE AND SEPARATOR MODIFICATION

• CELL EQUALIZATION

• STATE-OF-THE-ART - 140 Wh/kg

• PROBABILITY OF SUCCESS - 60%

C,F,_J%Z_SFRTI1

354
SPACEELECTROCHEMICAL RESEARCH& TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER A_DIL 11-1R, 1989

TECHNOLOGY ISSUES

Na - MCIx (160 Wh/kg)

• CELL REDESIGN

• CHEMISTRY CLARIFICATION

• TECHNOLOGY RESIDES OUTSIDE U.S,

• SEAL DEVELOPMENT

• TIME-DEPENDENT FAILURE MECHANISMS IDENTIFIED

• FREEZE/THAW IS REPORTED AS A NON-ISSUE

• PROBABILITY OF SUCCESS - 50°/_

G_J_A'_ FF_T I'_

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13, 1089

TECHNOLOGY ISSUES

ADVANCED Ni-H2 (71)Wh/kg)

• MODIFY SUPPORT STRUCTURES FOR COMPONENT, CELL

AND BATTERY

• TECHNOLOGY FLIGHT DEMO/QUALIFIED

• PROBABILITY OF SUCCESS - 95%

355
SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY

CONFERENCE

NASA LEWIS RESEARCH CENTER APRIL 11-13o 1080

TECHNOLOGY ISSUES

Li - SOLID CATHODE (200 Wh/kg)

• LI RECHARGEABILITY

• ELECTROLYTE STABILITY AND CONDUCTIVITY

• HIGHER ENERGY LONG-LIFE CATHODES REQUIRED (COO2)

• SCALE-UP FROM CONSUMER SIZES

• SAFETY

• PROBABILITY OF SUCCESS - 60%

C,_JSS_SERT04

SPACE ELECTROCHEMICAL RESEARCH & TECHNOLOGY CONFERENCE _,_

NASA LEWIS RESEARCH CENTER APRIL 11-13,1989

TECHNOLOGY ISSUES

LI - POLYMER ELECTROLYTE - SOLID CATHODE (200 Wh/kg)

• MANUFACTURING PROCESSING OF P.E.

• P.E. COMPOSITION AND CONDUCTIVITY

• INTERFACIAL CONTACT

• PROBABILITY oF SUCCESS . 60%

356

GH/JSg9S/RT 03

Z

w

=

=
Report Documentation Page National Aeronautics and

Space Administration

1. Report No. 2. Government Accession No. 3. Recipient's Catalog No.

NASA CP-3056

4. Title and Subtitle

Space Electrochemical Research and Technology (SERT) 1989

7. Author(s)

Richard S. Baldwin, Editor

9. Performing Organization Name and Address

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio 44135-3191

12. Sponsoring Agency Name and Address

National Aeronautics and Space Administration

Washington, D.C. 20546-0001

5. ReportDate

December 1989

6. Performing Organization Code

8. Performing Organization Report No.

E-4708

10. Work Unit No.

506-41-2 ]

11. Contract or Grant No.

13. Type of Report and Period Covered

Conference Publication

14. Sponsoring Agency Code

15. Supplementary Notes

16. Abstract

This document contains the proceedings of NASA's second Space Electrochemical Research and Technology

Conference, held at the NASA Lewis Research Center April 11-13, 1989. The objectives of the conference were

to examine current technologies, research efforts, and advanced ideas, and to identify technical barriers which

affect the advancement of electrochemical energy storage systems for space applications. The conference provided

a forum for the exchange of ideas and opinions of those actively involved in the field, with the intention of

coalescing views and findings into conclusions on progress in the field, prospects for future advances, areas

overlooked, and the directions of future efforts. Related overviews were presented in the areas of NASA

advanced mission models. Papers were presented and workshops were conducted in four technical areas:

advanced concepts, hydrogen-oxygen fuel cells and electrolyzers, the nickel electrode, and advanced rechargeable

batteries. This document contains the overviews, the technical papers, and summaries of the technical workshops.

17. Key Words (Suggested by Author(s))

Electrochemistry; Fuel cells; Batteries;

Electrocatalyst; Nickel electrode

18. Distribution Statement

Unclassified- Unlimited

Subject Category 44

19. Security Classif. (of this report) 20. Security Classif. (of this page) 21. No of pages

Unclassified Unclassified 364

NASAFORM1626OCT86 *For sale by the National Technical Information Service, Springfield, Virginia 22161

22. Price*

A16

NASA-Langley, 1989