UNITED STATES DEPARTMENT OF COMMERCE

C. R. Smith, Secretary

NATIONAL BUREAU OF STANDARDS • A. V. Astin, Director

Mass Transport In Oxides

Proceedings of a Symposium

Held at Gaithersburg, Maryland

October 22 - 25, 1967

Edited by

J. B. Wachtman, Jr. and A. D. Franklin

Institute for Materials Research

National Bureau of Standards

Washington, D.C. 20234

Sponsored by

Advanced Research Projects Agency

and

National Bureau of Standards

National Bureau of Standards Special Publication 296

Issued August 1968

For sale by ihe Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402

Price S3
Abstract

This Monograph presents the Proceedings of the Conference on Mass Transport in Oxides, held

at the National Bureau of Standards in Gaithersburg, Md., from October 22 to October 25, under the

joint sponsorship of the Bureau and the Advanced Research Projects Agency. Invited papers reviewed mass transport in alkali and silver hahdes as background, and topics relating to oxides including

point defects, (both experiment and theory), lattice dynamics and ionic interactions, atomistic mechanisms of diffusion and ionic conduction, chemical diffusion, defect complexes, assessments of the experimental situation for diffusion and ionic conduction, and the availability of good research materials. In contributed papers, attention was focused on color centers in alkaline earth oxides, cation

diffusion in MgO, transport of both cations and anions in transition metal oxides, and coulometric and emf techniques for obtaining transport data. A panel discussion summarized the work of the conference.

Key Words: Crystal defects, diffusion, ionic conductivity, mass transport, oxides, stoichiometry.

Library of Congress Catalog Card Number: 68-60030
Foreword

The National Bureau of Standards has as a major responsibility the task of ensuring that basic,

urgently needed data on the properties of materials are available to meet the requirements of the

Nation's scientists and engineers. In this context, data is a rather broad term, implying experimentally measured numbers, experimental procedures, theory, and the specification of the depend- ence of important properties upon controlling variables. In an immediate sense, this responsibility is met in part through measurements and calculations done in NBS laboratories, in part through compilation and pubhcation of critically evaluated

data gathered from the literature, and in part through survey articles including compiled data. To respond to the long-range needs, emphasis is placed on stimulating the development of fields where data will be increasingly needed in the future. Mass transport in oxides is an important factor controlling the properties and use of ceramic

materials at high temperatures and in controlling the processing of ceramics to achieve the properties needed for use at any temperature. Reliable data on diffusion coefficients and the factors controlling them are needed; however, this requires the development of reliable measurement

techniques, understanding of the relationship of structural defects to diffusion coefficients, and

high-quality, well-characterized research materials. In recognition of this need, the National

Bureau of Standards, through its Institute for Materials Research, and the Advanced Research

Projects Agency of the Department of Defense organized this Symposium on Mass Transport in Oxides, bringing scientists interested in diffusion together with those interested in such underlying fields as point defect studies and crystal growth. The intent was to provide a forum where

the problems currently limiting progress in obtaining rehable diffusion data might be more clearly

understood.

The National Bureau of Standards presents these Proceedings of the Conference in order to assist those attending the Conference in further research and to make the results available to a wider audience. On behalf of the recipient of these Proceedings, the National Bureau of Standards

would Uke to express appreciation to the participants, whose time, energy, and thought are embodied in these Proceedings, and to the Advanced Research Projects Agency, whose sponsorship

made the Conference possible.

A. V. Astin, Director

National Bureau of Standards.

Ill

Introduction

The Symposium on Mass Transport in Oxides was held at the National Bureau of Standards

laboratories at Gaithersburg, Maryland, on October 22-25, 1967. It brought together approximately

100 scientists and engineers from the United States, Great Britain, the Netherlands, Canada,

Austraha, and France. Mass transport in sohd, crystalhne oxides is an important factor controlhng the properties

and use of ceramic materials, directly when ceramics are used at high temperatures and indirectly

at any temperature, through influence on the processing and, therefore, the final properties of the ceramic.

Mass transport processes are directly important in high temperature apphcations involving creep, ionic conductivity, or protection from corrosion including oxidization. They are also directly

involved in fuel cell operation.

The indirect influence of mass transport processes arises through the dependence of such

structure-sensitive properties as strength, electrical and magnetic parameters, thermal conductivity, and optical properties upon the microstructure produced by processing. Most polycrystalline

ceramics are produced either by sohd state sintering or by a process involving partial or complete

melting foUowed by controlled crystaUization. Domination of the resulting microstructure is effected through careful control of the heat treatment (which may be rather complex) and through

the use of chemical additives to affect the rates of mass transport. An empirical approach has

generally been followed, sometimes quite successfully, but much effort is being devoted to improving ceramic processing, including the development of better quantitative descriptions in terms of the basic processes at work. An extensive survey of ceramic processing is given in the

report of the National Academy of Science-National Research Council Materials Advisory Board

ad hoc committee on Ceramic Processing, publication 1576 of the National Academy of Sciences.

Knowledge of mass transport rates and the factors controlling them is an area where improvement

is badly needed. Lack of reproducibility of diffusion data for oxides is a problem extensively dis- cussed in the present conference report and illustrated in a recent compilation of diffusion data

[A. L. Dragoo, Diffusion Rates in Inorganic Nuclear Materials, J. Res. NBS (Phys. and Chem.)

72A, 157 (1968)].

The field of mass transport in oxides at the present time contains much of importance and

opportunity to the basic science of materials, in addition to its technological significance. Detailed

mechanisms for similar processes in the "model" alkah hahdes have been rather well-worked out. While these models provide some guidance to understanding the behavior of oxides, they are

basically ideal-solution models and cannot be applied without perhaps very extensive modification

to the oxides, where defect concentrations are often much larger and where higher ionic charges

imply larger interactions among the defects themselves. It can be expected that considerable

progress in understanding the basic properties of soHds would result from more intensive study

of mass transport in oxides. The time appears ripe for expanding effort in this field, particularly

since recent advances in crystal growth techniques (particularly in the use of chemical vapor

transport) should now make it possible to prepare adequate research specimens.

There are- a number of detailed problems concerning mass transport in oxides that should be

considered. First, there is the question of the validity for oxides of the basic model developed for alkah hahdes, in which diffusion is described in terms of the concentration and jump frequencies

of isolated point defects. Evidently, means of identifying and measuring the concentration of dif- ferent species of point defects are important, as are means of calculating their properties. The

status of such calculations, especially for alkah hahdes, was recently surveyed in "Calculation

of the Properties of Vacancies and Interstitials," National Bureau of Standards Miscellaneous

Pubhcation 287, issued November 17, 1966. It is of interest to inquire how successfully the theories

described there apply to oxides. In particular, can the powerful new tools of lattice dynamics be

used to obtain the basic materials parameters needed for calculation of the properties of the defects

themselves? Theory taking into account details of the defect motion is then needed to express

the rate of mass transport in terms of point defect properties. Oxide structures are in many cases
more complex than the simple alkali halide structures; applications to these structures are es- sentially extensions of the traditional model within its conceptual framework. There is some

doubt about the adequacy of this framework, however. The effect of the partially-covalent bonding

in oxides upon defect properties is an open question and the existence of microdomains in many

oxides, which may be regarded as resulting from an ordering of defects, is a warning that the

traditional model may be too simple.

Second, how well do the phenomenological descriptions of chemical diffusion, developed

largely for alloys, describe behavior in oxides where in addition to the chemical composition

gradients directly associated with diffusing species there are effects associated with nonstoichiometry and with gradients of electrical potential?

Third, to what extent can ionic conductivity be separated from electronic conductivity and

used as a tool in the study of oxides in the same way that ionic conductivity has been so extensively

used in alkali hahdes? Conversely, can studies of electronic conductivity as a function of composition, environment, and temperature be used to infer the type and property of point defects and

to throw hght on the mechanisms of mass transport?

This conference was intended to provide a review of current understanding of mass transport

in oxides, and an opportunity to see where new research efforts might be fruitful. It consisted of a framework of invited papers chosen to cover most of the questions mentioned above, around which

were grouped contributed papers representing current work. The entire conference was summarized by a panel discussion, the report of which is included in these Proceedings.

The program began with a review of mass transport in alkali and silver halides and then turned

to the theoretical and experimental study of point defects in oxides, including the use of lattice dynamics in the investigation of interionic interactions. The atomistic mechanisms underlying the

phenomenological diffusion equations were then discussed, as well as the treatment of chemical

diffusion. The occurrence of defect complexes in oxides was treated and a general assessment was given of the state of existing data and the conditions needed to obtain reliable data. The

remaining invited papers dealt with ionic conductivity, coulometric studies, and the current

availability and the perfection of oxide single crystals.

No simple answers to the questions and problems previously mentioned were expected, but

suggested directions for work aimed at solving these problems were sought. The report of the

panel discussion summarizes the results. A few general points are mentioned here for emphasis:

1. The impurity level in the best currently available research specimens of oxides is high

enough that intrinsic behavior should not be expected in nearly stoichiometric oxides much below

the melting point. Better oxide single crystals are badly needed.

2. Departures from stoichiometry frequently occur in oxides. The most hkely possibiUty for

intrinsic and reproducible behavior is in systems where the nonstoichiometry is large enough to dominate impurity effects. Very careful control of thermodynamic variables and allowance of time

to equihbrate are necessary to fix the state of the specimen. This requirement holds for transport

work in oxides in general.

3. A difference of opinion exists concerning the emphasis which should be placed on detailed

study of atomic mechanisms of mass transport in simple and generally pure systems on the one hand and on phenomenological study of transport in engineering materials on the other hand. It appears essential to continue both, and probable that creative tension between proponents of the

two types of study wiU benefit both.

J. B. Wachtman, Jr. A. D. Franklin

Acknowledgment

The editors would hke to thank the following for their

help in organizing and conducting the Symposium:

C. E. Birchenall University of Delaware

R. E. Carter General Electric Company

J. H. Crawford, Jr. University of North Carohna

R. E. Howard National Bureau of Standards

J. R. Manning National Bureau of Standards

A. S. Nowick Columbia University

D. H. Whitmore Northwestern University

vi
Contents

Page Foreword HI Introduction V *Mass Transport in the Face-Centered Cubic Alkali and Silver Halides 1 L. M. SUfkin Diffusion of Sodium in Sodium Chloride in an Applied Electric Field 9

R. J. Friauf and V. C. Nelson *Characterization of Point Defects in Oxides 11

J. E. Wertz *Theory of the Energetics of Simple Defects in Oxides » 25

I. M. Boswarva and A. D. Franklin *Lattice Dynamics and Ionic Interactions 33

J. Slater On Electron-Lattice Interaction in the CaO F Center 39

J. C. Kemp and W. M. Ziniker Electronic States of Defects in Irradiated Oxides 41

B. Henderson Covalencv in Metal Oxides 51

B. E. F. Fender and B. C. Tofield *Diffusion and Ionic Conductivity: Kinetic Theory 53

J. R. Manning

*Chemical Diffusion Coefficients for Some Non-Stoichiometric Metal Oxides 65

J. B. Wagner, Jr. The Use and Limitations of the Concept of an Effective Binary Diffusion Coefficient for Multi-Component Diffusion 79

A. R. Cooper

Divalent Cation Impurity Diffusion in MgO 85

A. J. Mortlock Diffusion of Iron and Nickel in Magnesium Oxide Single Crysfals 89

S. L. Blank and J. A. Pask Impurity Cation Diffusion in Magnesium Oxide 95

B. J. Wuensch and T. Vasilos Cavity Formation in Magnesium Oxide 103

A. Briggs and D. H. Bowen

^Defect Complexes and Microdomains in Non-Stoichiometric Oxides, an Analysis of

Wustite, Fe,-60 109

G. G. Libowitz *Diffu8ion in Oxides: Assessment of Existing Data and Experimental Problems 119

C. E. Birchenall The Determination of Thermodynamic Properties in Single Phase Wiistite by

Coulometric Titration in a High Temperature Galvanic Cell 129

H. F. Rizzo. R. S. Gordon, and I. B. Cutler The Mechanism of Oxygen Self-Diffusion in Nickel and Cobalt Oxides 143 M. Hoch and R. Szwarc Experimental Evidence for Highly Mobile Electrons in MnO and NiO at High

Temperature 145

D. S. Tannhauser, N. M. Tallan, and M. Gvishi High Temperature Defect Structure and Electrical Properties of NiO 147

I. Bransky and N. M. Tallan *Ionic Conductivity in Oxides: Experimental Problems; Survey of Existing Data 149

L. Heyne

Measurement of High-Temperature Thermodynamic Properties of Non-Stoichiometric Oxides Using Solid State EMF and Coulometric Techniques 165

B. C. H. Steele Interdiffusion Coefficients From Electronic Conductivity Measurements — Application to CuiO 173 R. H. Campbell, W. J. Kass and M. CKeeffe

*Invited Paper.

Vll
Page A Chemla Experiment in BeO 177

C. F. Cline, H. W. Newkirk, R. H. Condit, and Y. Hashimoto

Self-Diffusion of Oxygen in Neodymium and Samarium Sesquioxide 179 G: D. Stone, G. R. Weber, and L. Eyring

Oxygen Transport During Oxidization 187

J. B. Lightstone and J. P. Pemsler Mechanical and Dielectric Relaxation of Hopping Electrons and the Ionic Defects

in Reduced Rutile (TiO.-x) 189 W. W. Scott and R. K. MacCrone

Purity and Perfection of Research Specimens of Oxides 195

J. W. Cleland The Growth of Oxide Single Crystals by Chemical Transport 205 Robert Kershaw and Aaron Wold

PANEL DISCUSSION 207

A. S. Nowick (Chairman), R. E. Carter, J. H. Crawford, Jr., Y. Haven, A. B. Lidiard,

J. B. Wachtman, Jr., D. H. Whitmore.

Vlll
Mass Transport in the Face-Centered Cubic Alkali and Silver Halides*

Lawrence Slifkin

Department of Physics, University of North Carolina, Chapel Hill, N.C. 27515

Some current concepts and problems relating to mass transport in alkali and silver halide crystals are discussed: (1) There are contributions from several types of defects to ionic conductivity and

tracer diffusion; thus, the analysis of such data to yield energies and entropies of formation and migration is not a trivial matter. (2) In a few cases, the energy and entropy of formation of the defect pair have been resolved into the components associated with each member of the pair. (3) Bound Schottky

pairs have been identified, and in a few materials their characteristics have been measured. (4) The

diffusion of impurity ions in ionic crystals still presents some unanswered questions. (5) Several other problem areas are listed, the study of which thus far is in a relatively early stage.

Key Words: Alkali hahdes, diffusion, ionic conductivity, silver halides, tracer diffusion.

1. Why Halide Crystals?

Our present understanding of mass transport in oxide crystals is very limited indeed. Not only are pure crystals difficult to prepare, but the melting

points of many of these materials are much higher

than the range of the usual high-temperature fur- naces. As a result, although there have been a number of experiments on diffusion in the extrinsic,

impurity-sensitive range, one can only speculate as

to the nature of intrinsic defect phenomena.

In contrast, the alkali haUdes have readily ac- cessible melting points and may be prepared as

crystals of very high purity; studies of the intrinsic mass transport properties are thus quite feasible. During the past 20 years there has been an intensive

investigation of these materials [2, 67],' mainly by means of measurements of ionic conductivity and

tracer diffusion, but also employing dielectric loss and magnetic resonance techniques. The materials

which have been examined in most detail are NaCl,

KCl, and KBr, aU exhibiting Schottky defects; and

AgCl and AgBr, in which Frenkel defects predominate. Our understanding of these materials, while

certainly not flawless, has reached a level that is not hkely to be attained for the oxides for some time

to come. It thus seems reasonable that the concepts,

problems, and difficulties that have emerged from

the study of these halides may well provide a useful

guide in the development of a science of ionic trans- port phenomena in oxide crystals.

2. Ionic Diffusion and Conductivity

Let us consider, for concreteness, a crystal containing Schottky defects (dissociated cation and

•Preparation of this manuscript supported by the U.S. Atomic Energy Commission

(Contract No. AT(40-1)-2036), the Office of Aerospace Research (Grant No. AFAFOSR-450-66 ) and the Advanced Research Projects Agency (Contract No. SD-100 )

.

* Figures in brackets indicate the hterature references at the end of this paper.

anion vacancies), along with a small concentration

of divalent metal ion impurity. The diffusion coefficient of either type of vacancy is given by [1, 2]

o

Here, A. is the jump distance, presumably the dis- tance to the nearest equivalent site; v is the number

of jumps made per second by the vacancy, and is equal to

p=12vo exp [-(H,n-TS,n)lkT].

The factor 12 enters because there are 12 equivalent

neighbors in the face-centered cubic lattice; vo is an attempt frequency, usually taken to be the Debye

frequency because of lack of a more exact description; Hm is the enthalpy (or, for most purposes at atmospheric pressure, an energy) of activation for migration; and the entropy of migration, Sm, results from transient changes in lattice mode frequencies

during the jump process. Motion of these vacancies produces self-diffusion of the corresponding ion, with a coefficient given by

where /, the correlation factor, allows for the fact that the tracer does not make successive jumps at random, even though the vacancies do (for f.c.c. lattices, /= 0.78); x is the fractional concentration

of this type of vacancy (i.e., either cation or anion

vacancy). The concentrations of anion and cation

vacancies satisfy a relation of the mass-action type:

xjX2 = exp • [—(Hf—TSf)lkT],

the subscript/ referring to the process of formation

1
of the dissociated Schottky pair of vacancies. Thus,

if the crystal is sufficiently pure and the temperature sufficiently great that the concentrations of

impurities are negligible compared to Xi and X2, so that the two types of vacancy are present in equal

concentrations, then either concentration increases

with increasing temperature as exp {—Hfl2kT).

On the other hand, if the concentration of, say, divalent cation impurity is not neghgible, then the

cation vacancy concentration is increased and the

anion vacancy concentration is suppressed. Al- though each such impurity ion must, because of charge neutrality, introduce a cation vacancv, not

all of these extra vacancies will be free. There will be at least a coulombic attraction between impurity

and vacancy, and in the simplest model we may

consider the two to form a new species, a "complex," when they are nearest neighbors, and to be

free otherwise. This equilibrium is similarly de- scribed by a mass-action equation.

=r2-exp[{Ha-TSa)lkT],

where Xk is the fractional concentration of complexes, c is the total fractional concentration of divalent cation impurity, and Ha and Sa are the

enthalpy and entropy given up in the association process. In addition, electrical neutrality of the

crystal requires that X1—X2+ (c — Xk)

Ionic mobility is more frequently studied by means of the resulting electrical conductivity than

by self-diffusion experiments. If charge and the

tracer are both transported by the same mechanism, then the Nernst-Einstein equation leads to the following relation between D and conductivity, cr:

where is the number of molecules per unit vol- ume and q is the magnitude of the charge trans- ported by each carrier. From these relations, and on the simple association model which we have considered, it is thus

possible to calculate as a function of c and T, the

contribution from each defect to the ionic conductivity and self-diffusion — provided one knows the

values of the three sets of enthalpies and entropies

introduced above. In practice, of course, we wish

to invert the calculation; we know the total conductivity and/or various diffusion coefficients, and from

this it is necessary to extract the various enthalpies

(essentially energies) and entropies for each defect.

3. Defect Formation and Migration

Energies

If the concentration of impurities is small and

if only one type of defect, say the cation vacancy,

contributes to the conductivity, then the Arrhenius

plot of log (err) versus 1/7" would be a straight line

at high temperatures, with a slope equal to — (f Hf+ Hm) Ik. At lower temperatures, where the

concentration of intrinsic, thermally induced defects

falls well below that of divalent cation impurity,

then if one could neglect the impurity-vacancy

association the slope should be equal to —HnJk.

At still lower temperatures, the formation of complexes becomes significant, and the slope would

increase again, reaching the value — (//,«+ | Ha) Ik.

It should therefore be possible to determine the

values of these three energies from such an Arrhenius plot of the conductivity, and this has indeed

been done for many of the hahdes [3, 4]. Typical

values for the common alkali halides are: 2 to 2V2 eV for the formation of a Schottky defect,

0.6 to 0.8 eV for the activation energy for migration

of the cation vacancy, and 0.3 to 0.5 eV for the

impurity-vacancy association energy. Unfortunately,

however, the expected linear region below the in- trinsic range and above the association region is often so short that there may be considerable ambiguity in determining the slope. Various studies

of the same material have thus occasionally resulted

in values of //„, differing by about 0.1 eV, with

comparable uncertainties produced in the other

energies. Moreover, it is difficult to be certain that only one type of defect is moving in the high temperature region.

An alternative approach has been to examine the

variation of ionic conductivity with increasing impurity concentration. The analysis of such conductivity isotherms is based on manipulation of the

several equations given above, and has been lucidly

formulated by Lidiard [2]. Generally, certain

approximations must be made, such as the assumption that single cation vacancies are solely responsible for the conductivity — at least in the lower

temperature region of the intrinsic range — or that the entropy of the impurity-vacancy association can be taken as zero. Moreover, the simple model

which we have considered does not adequately

account for extended, or "excited," configurations

of the complexes nor for long-range electrostatic

screening effects of the Debye-Hiickel type, which

Lidiard [2] has shown to be important at high

concentrations of point defects. For the silver halides, self-diffusion experiments

[5, 6| have conclusively demonstrated that anion

transport is negligible in comparison with cation

mobility, but the situation is complicated by a marked upward curvature in the Arrhenius plots
at high temperatures, a curvature that has been

interpreted in terms of Debye-Hiickel interactions

[7] and also partially as an effect of an excess lattice expansion [8]. There are two recent extensive

studies of silver chloride, Abbink and Martin's

analysis of conductivity isotherms [7] and Miiller's

[9] analysis of the temperature-dependence of the

conductivity in the region of the "knee" between

intrinsic and extrinsic ranges; the results from the

two differ considerably. Abbink and Martin per- formed a careful experiment, but the results of

their analysis are sensitive to the accuracy of the

Debye-Hiickel treatment. The resolution of this conflict is certainly not obvious.

Additional information on silver ion transport in AgCl and AgBr has been gleaned from detailed and

precise comparisons of ionic conductivity with the

silver tracer diffusion coefficient, looking at apparent

deviations from the Nernst-Einstein equation [6, 10, 11, 12]. The main point of interest here is that

in an "interstitialcy" jump, where the interstitial ion replaces an ion in an adjacent normal site, the

jump distance for the electrical charge is greater

than that for the radioactivity. Recent work of

Friauf and co-workers [11, 12] has actually been

able to specify the contributions from coUinear and

from noncoUinear interstitialcy jumps. Studies of

halide self-diffusion in these materials, however,

have not yet given such a complete picture. In both the chloride [13] and the bromide [14], the

Arrhenius plots are quite curved at low temperatures, suggestive of several alternative diffusion

mechanisms.

The alkali halides present somewhat different problems. Since the concentrations of thermal

defects here are appreciably smaller than in the

silver haUdes, Debye-Hiichel screening effects are

not expected to be as important (although Kanzaki

[15] has interpreted deviations from predictions of

the simple theory of the conductivity isotherms for calcium-doped NaCl as due to long-range electrostatic interactions). On the other hand, a contribution to the ionic conductivity from anion vacancies

poses a worrisome problem. This sometimes appears

as a curvature in the intrinsic region of the Arrhenius plot [16, 17], but a straight line is no guarantee

of a single process. Thus, in KBr [18], although

both types of vacancy make comparable contributions, and the activation energy for the anion

vacancy is 0.2 eV larger than that for the cation

vacancy, nevertheless, the intrinsic region of the

Arrhenius plot appears to be quite linear. Another

difficulty that plagues the analysis of conductivity

isotherms is accurate knowledge of the effective

impurity concentration. This is not only a problem

of measuring the magnitudes of deliberate additions,

but also of determining the presence of background

impurities which can react with a portion of the

added dopant; e.g., small quantities of 0H~, which

may react with added Sr++ [17]. How, then, must one deal with the alkali halides?

One approach applied by Rolfe [18] to KBr, is to determine conductivity isotherms of crystals doped

with divalent anions as well as with divalent cations.

This procedure, however, has not been generally

extended because no other f.c.c. alkali halide has

been found to offer an appreciable solubility for a divalent anion. A second recent approach is to instruct a computer to obtain those values of the

various energies and entropies which will provide

the best fit to carefully measured Arrhenius plots for various concentrations of impurity [19, 20, 21].

There are two energies and entropies of migration

(one set for each type of vacancy), along with corre- sponding parameters involving Schottky defect

formation and impurity-vacancy binding (the binding

entropy is sometimes hopefully set equal to zero).

With so many adjustable parameters to be employed

in fitting a rather unspectacular, monotonic Arrhenius plot, it is not at all obvious that the procedure

should work. Nonetheless, unique values do emerge,

values which are consistent among Arrhenius plots

involving different impurity concentrations. The

method is not without its hazards, however, as Dawson and Barr [21] have shown that the best set of values may be rather sensitive to the extent of

the temperature range used in fitting theoretical to experimental conductivities.

By far the most rehable and convincing procedure

would seem to be a combination of ionic conductivity measurements with tracer studies of anion

and cation self-diffusion, as has been done recently

by Fuller and Reilly [20] and by Dawson and Barr

[21]. Diffusion experiments have the advantage

that they are not influenced by mobility of vacancies

of the opposite type. On the other hand, mobile

neutral species, such as impurity-vacancy complexes and bound vacancy pairs, which would be

invisible in a conductivity measurement, may make

an appreciable contribution to tracer diffusion. The

importance of the impurity-vacancy complex may

be minimized by working at high temperatures; the

contribution from bound vacancy pairs must be

separately evaluated. This is accomplished by ob- serving the effect of additions of divalent cation on

the diffusion coeflicient of the anion. The cation

impurity adds cation vacancies, thereby suppressing

the anion vacancies, but leaves unchanged the

concentration of bound pairs [20, 21, 22, 23]. In

particular, Dawson and Barr [21] have recently

obtained excellent agreement between results of tracer diffusion studies in KBr and computer fits of Arrhenius plots.

The end products of such experiments are values

3
of ///, Hm for both anion and cation vacancy. Ha

for complexes involving various impurities, and all of the corresponding entropies. Perhaps one of the

most valuable uses of the enthalpies is to provide

a guide for the theoretician: the interionic potentials

and polarization corrections in these materials are somewhat uncertain, and trial functions can be

tested by comparing the predicted energies with

the experimental ones. The entropies, however,

presently are by-products for which little use has

been found. One can qualitatively understand that these entropies are so much larger than in metals

because of the large relaxations which occur when

an ion is removed from an electrostatically bound

structure; a quantitative theory, unfortunately, is

still lacking. It is also becoming apparent from

recent work that anion vacancies and bound pairs

often make a significantly larger contribution to transport phenomena than had previously been

appreciated. Lidiard [24J has further pointed out

that the defect formed by addition of a vacancy to a bound pair should have some stability and

mobility, and may make a nontrivial contribution.

4. Resolution of Hj and Sf into Their

Components

The separate components of a Schottky or Frenkel pair can be individually created or an- nihilated at certain jogs on dislocations and surfaces.

The dislocation, for example, must therefore be in equiHbrium separately with both component point

defects, and this will usually leave it with a net charge [25J. If there is an excess of cation vacancies

in the crystal, because of the presence of divalent

cation impurities, some of these will condense onto

the dislocation, giving it a negative charge. At high

temperatures, on the other hand, one can neglect

the effects of impurities and if the free energy of formation of the cation vacancy is less than that

for the anion vacancy, the dislocation will produce

an excess of cation vacancies and will acquire a

positive charge. At some intermediate temperature,

the isoelectric point, the dislocation is electrically neutral; from the dependence of this temperature

upon impurity concentration it is possible to deduce

the energy and entropy of formation of the cation

vacancy. These quantities may then be subtracted

from those for the complete Schottky or Frenkel

pair to give the parameters characterizing the other

component of the pair.

Kanzaki, Kido, and Ninomiya [26J studied the

effects of dislocation density on the ionic conductivity of very pure and doped KCl. They interpreted

the results in terms of the effect of the compensating

space charge surrounding the charged dislocations,

and were able thereby to deduce the vacancy formation free energy at one particular temperature.

Davidge [27J, Strumane and DeBatist [28], and

Kliewer and Koehler [29J determined the concentration dependence of the isoelectric point in NaCl

by observation of potential differences produced

by dislocation displacement and by studies of the

effect of dislocations on the elastic modulus. Their

results are essentially in agreement with one

another and indicate that the formation energy of

the cation vacancy is only shghtly less than that

for the anion vacancy, but that the formation entropy

for the cation vacancy is considerably larger than

that for the anion vacancy. Flint and Breig [30] have

deduced vacancy formation free energies from light

scattering, but for NaCl the result does not agree

with Kliewer, etc. Charged dislocation effects in silver chloride

have been calculated by Kliewer [31]. He shows

how the vacancy formation entropy can be obtained

from Wakabayasi's data [32] on the temperature

dependence of contact potential and he gives an

estimate for the energy of formation. This energy

has recently been measured [33], and the probable

effects of the surface charge on the mechanism of

the photographic process are cited. As in sodium

chloride, the cation vacancy has a smaller formation energy and a larger formation entropy than

does its complementary defect (the interstitial silver ion, in this case). Saunders, Tyler, and West

[34] have been able to probe the potential near the

surface of AgBr crystals by an ingenious experiment in which photographic development tech- niques are exploited to foUow the drift of photoelectrons in thin specimens. They find that at room

temperature the surface is negatively charged,

as in AgCl [33], but their results are interpreted

by Trautweiler [35] in terms of intrinsic, rather

than impurity effects. This seems somewhat questionable, since it would then be necessary that a

purity of 0.1 ppm be preserved in AgBr slabs only

a few tens of microns thick, formed by melting

between glass plates.

5. Bound Schottky Pairs

Effects of bound Schottky pairs in ionic crystals were perhaps first seen by Tannhauser [5] in 1958,

who was puzzled by curvature extending over the

entire temperature range of the Arrhenius plot

for diffusion of bromide in silver bromide. In various alkali halides, bound pair contributions to anion diffusion were found by Morrison and co- workers [36], Laurance [22], and Fuller [20]. The

vacancy pair diffusion contribution in KBr has re- cently been measured by Barr and Dawson [14]. Downing and Friauf [37] compared the ionic conductivity in NaCl with the diffusion of sodium tracer

4
and have found an "anomalous" contribution to diffusion which is attributed to bound p^irs.

Dielectric loss experiments in several of the

alkah halides by Sastry, Srinivasan, and Economou

[38] have shown peaks which have many of the

characteristics of vacancy pair relaxation processes. The pair formation energies, rotation activation energies, and rotation preexponential factors are aU consistent vrith what would be expected for

pairs. Doping with divalent cation impurity, which

would perturb the monovacancy concentrations

and the conductivity but not the concentration of

pairs, was found not to affect the results. Also,

measurements on crystals of different thicknesses seem to show that the effect is a bulk one and not a surface phenomenon of the type studied by Wimmer

and Tallan [39]. The one disturbing fact, and it is quite disturbing indeed, is that the intensities of

the losses are larger than expected, by a factor of

about 50 to 100. Boswarva and Franklin [40] have

concluded that extending the simple Debye theory

to include long range interactions cannot patch

up the discrepancy. Recently, Chandra and Agrawal

[41] have calculated the entropies of formation of

bound pairs in alkali hahdes, have found them

to be quite large, and have concluded that the very

large concentrations of pairs implied by the ex- periment are reaUy correct. This leads to predictions, however, of startHngly large density anomalies, specific heat anomalies, discrepancies

between diffusion and conductivity, etc., in conflict with the lack of experimental observation of such

anomalies. Sack and Smith [41a] have observed,

in several halides, an extra dielectric loss at high

temperatures and at microwave frequencies which

is attributed to relaxation by vacancy pairs and

which does not show an anomalously large in- tensity. It thus seems rather likely that the large

concentrations of vacancy pairs deduced by

Sastry et al. and by Chandra and Agrawal are erroneous.

If the rotational activation energy for vacancy

pairs could indeed be convincingly determined

from a relaxation experiment of this sort, a comparison could be made with the pair migration

energy obtained from diffusion studies. The two

energies are not the same. Of the two types of jump

which a bound vacancy pair can make, the relaxation experiment would measure the faster, while

the diffusion experiment is Umited by the slower.

If both activation energies could thus be obtained,

a comparison with the predictions of Tharmalingam and Lidiard [23] would be of interest. Vacancy pairs were studied in AgCl by Layer

[42], by quenching very pure material from temperatures near to the melting point. Correction of

the original data for more recent values of the

vacancy migration energy, etc., yields a binding

energy of the pair of 0.5 eV, a migration activation

energy of 1.0 eV, and a quenched in concentration

of approximately 10"''. This concentration pre- sumably refers to dissociated Schottky pairs at the

original quench temperatures. It is sufficiently

small not to contradict the nonobservance of Schottky defects in AgCl by mass-transport studies

[12] or combined measurements of lattice parameter

and density [43]. The results are consistent with

self-diffusion experiments [6, 13] and with the observation of very large dislocation loops in nearly

perfect regions of AgCl crystals which appear to have been produced by condensation of Schottky

defects [44]. 6. Impurity Ion Diifusion

The diffusion of impurity ions is considerably more complicated than is self-diffusion. Except

for a measurement by Batra [45] of the diffusion of

bromide and iodide tracers in silver chloride (in which it was — surprisingly — found that they both

have the same activation energy as does the

chloride ion), there has been virtually no recent work on the diffusion of anionic impurities. The

remainder of this discussion is thus concerned

only with cation diffusion.

It is generally assumed that metallic ions in alkah hahde crystals simply substitute for the

alkali ions. That this may not always be true is indicated by the optical absorption spectrum of

Co++ which in several halides is that characteristic

of the symmetry of an interstitial, not substitional,

site [46, 47]. There is also some evidence from

electrolysis experiments [48] that Pb++ forms a negatively charged complex in KCl, or at least in KCl with certain impurities present. In the silver halides, in view of the well-known ease with which

the silver ion enters interstitial sites, one is not surprised that Cu+ [49] and Au+ [50] also diffuse

interstitially. It is, however, somewhat unexpected

that Fe+ + + resides interstitially in AgCl in a complex involving, at least at low temperatures, a tetrahedron of silver vacancies [51, 52].

It has been emphasized by Lidiard [2] that for the usual, substitutional case, the impurity ion can

only migrate when it is part of an impurity-vacancy

complex. Also since a polyvalent ion affects the

local concentration of vacancies, the diffusion

behavior depends very much on whether the concentration of the diffusing ion is large or small as compared to the "background" concentration of vacancies. This discussion has been extended

recently by Howard and Lidiard [53]. The treatment of divalent cation impurity diffu- sion centers about the properties of the impurity5
vacancy complexes. These have been studied by

a variety of techniques: ionic conductivity, di- electric loss, depolarization effects, anelasticity,

electron paramagnetic resonance, and tracer

diffusion. Reported binding energies are generally

in the range 0.3 to 0.5 eV in the alkaU hahdes and

somewhat less in the silver halides, but there is Uttle agreement within this range among the various

experiments on a given system. Watkins [54] has

shown, by means of magnetic resonance and di- electric loss experiments, that in Mn-doped NaCl

the binding energy at the next-nearest-neighbor site

is almost as large as at the nearest-neighbor site.

It is also known that larger complexes may be

formed by aggregation of simple impurity-vacancy

pairs, the initial step being production of a trimer

[55].The cation vacancy in the neighborhood of a metallic impurity ion has a number of different jump probablilites: it may exchange with the impurity, rotate about the impurity, dissociate from

the impurity, and jump from a more distant site to a site adjacent to the impurity. The ratio of these

last two, of course, is determined by the binding

free energy. The frequencies of all four types of

jump, however, must be known if a complete understanding of impurity diffusion is to be attained. It

is expected that the extra charge on the divalent

impurity ion would increase the activation energy

for its exchange with an adjacent vacancy, and

experiments by Dreyfus and Laibowitz [56] on

anelasticity and dielectric loss as a function of

crystal orientation do show that the vacancyimpurity exchange frequency is less than that for the jumping of the vacancy around the impurity. The

complex is thus able to rotate about the impurity

much faster than it can actually diffuse. Further

details about these jump frequencies are lacking.

With the analysis of impurity diffusion in terms

of the various jump frequencies presently impossible, recent work on impurity diffusion in haUde

crystals has been concerned with the dependence

of the diffusion coefficient on impurity concentration, as in the experiments of Keneshea and Fredericks [57] and of Rothman, Barr, Rowe, and

Selwood [58] on diffusion in KCl and NaCl. In

general, most of these results can be described

by the simple association theory of Lidiard. There

are available three studies of divalent cation dif- fusion in silver haUdes. Hanlon [59] measured the

diffusion of cadmium in pure and in cadmiumdoped AgBr, and found that the data required

the assumption that the vacancy mobility increased

by a factor of two for each 1 atomic percent cadmium present. Reade and Martin [60] diffused

cadmium tracer into pure AgCl; Laskar [61] dif- fused manganese tracer into pure AgCl. In all

three cases of tracer diffusion in pure silver halide,

the Arrhenius plots appear to consist of two seg- ments, with a break at approximately 325 °C. It is possible that the tracer may be diffusing by both

vacancy and interstitial mechanisms, but at present

this is at best only speculation.

7. Several Other Problem Areas

In addition to those problems which have been

discussed in the preceding sections, there are several others which are of considerable interest. Most of these are in rather early stages of de- velopment.

(a) Electrode polarization effects [62] have been

studied both experimentally and theoretically,

but there are still wide gaps between the two.

(b) Effects of high pressure, described by acti- vation volumes, have been measured in only a few

cases [4, 63].

(c) Thermoelectric effects have been discussed

by Howard and Lidiard [53]. Most of the data that are available are for the silver hahdes; Christy's

results [64] on AgCl indicate that the heats of trans- port are quite large and are temperature dependent.

(d) Little is known of the diffusion along dislo- cations in ionic crystals; how wide are the "pipes"?

what are the activation energies? what are the ef- fects of the charge on the dislocations? what effects arise from impurity segregation? [66].

(3) Llity [65] has measured the migration energy

for the anion vacancy in KCl by means of color center reactions at low temperatures. The experiments quite conclusively give 0.6 eV, or only ap- proximately 2/3 the value obtained at high temperatures from tracer diffusion and ionic conductivity

experiments. If the ionic configuration around the

vacancy does indeed change at some intermediate

temperature, so that the two experiments are observing different kinds of anion vacancy, then a new dimension will have been introduced into the

study of mass transport in ionic crystals.

8. Concluding Remark

The alkali and silver hahdes thus stiU pose a few

mysteries and many unanswered questions, but

these are, after all, quite detailed and intimate

questions. It would appear that we now have both

a sound conceptual base and also the techniques

necessary for obtaining the answers. This is a

position which justifies cautious optimism; it has

been reached, however, only as a result of interpretation of measurements extending over large ranges

of temperature and purity, penetrating well into both the intrinsic and extrinsic regions. Had one

6
been limited to temperatures below half of the melt- ing point and to purities poorer than a hundred

parts per million, as has so far been the case in research on oxides, then it is doubtful whether

many of the present concepts could have been de- veloped, or whether we would even know what

questions to ask. The road ahead for the oxides thus

appears to be long and difficult indeed.

9. References

[1] See, for example, P. Shewmon, Diffusion in Solids (McGraw-Hill Book Co., New York, N.Y., 1963).

[2] A particularly useful discussion of mass transport in ionic crystals is: A. Lidiard in Handbuch der Physik, XX,

p. 246, ed. by S. Fliigge (Springer-Verlag, Berlin, 1957).

[3] See, for example. R. Dreyfus and A. Nowick, J. Appl.

Phys. 33 , 473 (1966); S. Jain and S. Dahake. Phys. Letters 3 , 308 (1963); Indian J. Pure & Appl. Phys. 2, 71

(1964).

[4] A. Abey and C. T. Tomizuka, J. Phys. Chem. Solids 27,

1149 (1966).

[5] D. Tannhauser, J. Phys. Chem. Solids 5, 224 (1958).

f6J W. D. Compton and R. Maurer, J. Phys. Chem. Solids 1, 191 (1956).

[7] H. Abbink and D. Martin, Jr., J. Phys. Chem. Solids 27,

205 (1966).

[8J P. MiiUer, Phys. Stat. Sol. 21, 693 (1967).

[9] P. MuUer, Phys. Stat. Sol. 12, 775 (1965).

[10] A. Miller and R. Maurer, J. Phys. Chem. Solids 4, 196 (1958).

[11] R. Friauf, J. Phys. Chem. Solids 18, 203 (1961); J. Appl.

Phys. 33, 494 (1962).

[12] M. Weber and R. Friauf, to be published.

[13] E. Lakatos and K. Lieser, Z. Physik. Chem. 48, 213 (1966).

[14] L. Barr and D. K. Dawson, private communication.

[15] H. Kanzaki, K. Kido, S. Tamura, and S. Oki, J. Phys. Soc. Japan 20 , 2305 (1965).

[16] A. AUnatt and P. Jacobs, Trans. Faraday Soc. 58, 116 (1958).

[17] P. Jacobs and J. Maycock, Jr., J. Phys. Chem. Solids 24,

1693 (1963).

[18] J. Rolfe, Can. J. Phys. 42 , 2195 (1964).

[19] J. Beaumont and P. Jacobs, J. Chem. Phys. 45, 1496 (1966).

[20] R. Fuller, Phys. Rev. 142 , 524 (1966); R. Fuller and M.

ReiUy, Phys. Rev. Letters 19, 113 (1967).

[21] D. K. Dawson and L. Barr, Phys. Rev. LeUers 19, 844 (1967).

[22] N. Laurance, Phys. Rev. 120, 57 (1960).

[23] K. Tharmalingam and A. Lidiard, Phil. Mag. 6, 1159 (1961).

[24] A. Lidiard, oral comments, this conference.

[25] J. Frenkel, Kinetic Theory of Liquids (Oxford University

Press, 1946); K. Lehovec, J. Chem. Phys. 21, 1123 (1953).

[26] H. Kanzaki, K. Kido, and T. Ninomiya, J. Appl. Phys. 33,

482 (1962).

[27] R. Davidge, Phys. Stat. Solidi 3, 1851 (1963).

[28] R. Strumane and R. DeBatist, Phys. Stat. SoHdi 6, 817

(1964).

[29] K. Kliewer and J. Koehler, Phys. Rev. 157, 685 (1967).

[30] C. PUnt and M. Breig, J. Appl. Phys. 35 , 2745 (1964).

[31] K. KUewer, J. Phys. Chem. SoUds 27, 705, 719 (1966).

[32] H. Wakabayasi, J. Phys. Soc. Japan 15, 2000 (1960).

[33] L. Slifkin, W. McGowan, A. Fukai, and J. Kim, Phot. Sci. & Eng. 11, 79 (1967); J. Kim, Ph.D. Thesis, University

of N.C., 1967; W. McGowan, Ph.D. Thesis, University

of N.C., 1965.

[34] V. Saunders, R. Tyler, and W. West, to be published.

[35] F. Trautweiler, to be published.

[36] J. Morrison and R. Rudham, J. Phys. Chem. Solids 6, 402

(1958); L. Barr, J. Morrison, and P. Schroeder, J. Appl.

Phys. 36 , 624 (1965).

[37] H. Downing and R. Friauf, private communication.

[38] P. Sastry and T. Srinivasan, Phys. Rev. 132, 2445 (1963); N. Economou, Phys. Rev. 135, A1020 (1964); N. Econ- omou and P. Sastry, Phys. Stat. Solidi 6, 135 (1964).

[39] J. Wimmer and N. TaUan, J. Appl. Phys. 37, 3728 (1966).

[40] I. Boswarva and A. Franklin, Phil. Mag. 11, 335 (1965).

[41] S. Chandra and V. Agrawal, J. Phys. Chem. Solids 28,

1055 (1967). [41a] H. Sack and G. Smith, J. Phys. Soc. Japan 18, Suppl. 3, 124 (1963).

[42] H. Layer, M. Miller, and L. Slifkin, J. Appl. Phys. 33, 478

(1962). [43] R. Fouchaux and R. Simmons, Phys. Rev. 136A, 1664

(1964).

[44] C. Childs, A. Fukai, W. McGowan, and L. Slifkin, to be

published.

[45] A. Batra and L. Slifkin, to be pubhshed; see also J. Laurent and J. Benard, J. Phys. Chem. SoUds 3, 7 (1957).

[46] M. Hills, J. Phys. Soc. Japan 19 , 760 (1964); K. Lai, Phys.

Stat. Solidi 22, 141 (1967).

[47] M. Musa, Phys. Stat. Sohdi 16, 771 (1966).

[48] W. Fredericks and A. Scott, J. Chem. Phys. 28, 249 (1958).

[49] P. SUptitiz, Phys. Stat. Solidi 7, 653, 667 (1964).

[50] A. Batra, A. Laskar, and L. Slifkin, to be published.

[51] W. Hayes, J. Pilbrow, and L. Slifkin, J. Phys. Chem. Solids 25, 1417 (1964).

[52] M. Satoh and C. Slichter, Phys. Rev. 144, 259 (1966).

[53] R. Howard and A. Lidiard, Rept. Prog. Phys. XXVII, 161

(1964).

[54] G. Watkins, Phys. Rev. 113, 79, 91 (1959).

[55] J. Cook and J. Dryden, Proc. Phys. Soc. 80, 479 (1962); Y. Chiba, K. Veki, and M. Sakamoto, J. Phys. Soc. Japan

18, 1092 (1963); H. Symmons and R. Kemp, Brit. J. Appl.

Phys. 17, 607 (1966); K. Shrivastava, Nuovo Cimento

47B, 251 (1967).

[56] R. Dreyfus and R. Laibowitz, Phys. Rev. 135, A1413 (1964).

[57] F. Keneshea and W. Fredericks, J. Chem. Phys. 38, 1952

(1963) ; 41, 3271 (1964); 43, 2925 (1965); 46, 2000(1967);

J. Phys. Chem. Solids 26, 501, 1787 (1965).

[58] S. Rothman, L. Barr, A. Rowe, and P. Selwood, Phil. Mag.

14, 501 (1966).

[59] J. Hanlon, J. Chem. Phys. 32, 1492 (1960).

[60] R. Reade and D. Martin, Jr., J. Appl. Phys. 31. 1965 (1960).

[61] A. LaskarandL. Slifkin, Bull. Am. Phys. Soc. 11,838(1966).

[62] For recent work, see A. Allnatt, P. Jacobs, and J. Maycock,

J. Chem. Phys. 43 , 2526 (1965); 44, 3644 (1966); J. Beaumont and P. Jacobs, J. Phys. Chem. Solids 28,

657 (1967); H. Wintle and J. Rolfe, Can. J. Phys. 44, 965

(1966); 45, 2253 (1967); D. Kahn and J. Maycock, J. Chem.

Phys. 46, 4434 (1967).

[63] S. Kurnick, J. Chem. Phys. 20, 218 (1952); W. Taylor, W. Daniels, B. Royce, and R. Smoluchowski, J. Phys

Chem. Solids 27, 39, 1196 (1966); C. B. Pierce, Phys

Rev. 123, 744 (1961).

[64] R. Christy, J. Chem. Phys. 34, 1148 (1961): R. Christy, Y. Hsueh, and R. Mueller, J. Chem. Phys. 38, 1647 (1963); see also P. Jacobs and J. Maycock, Trans. AIME 236,

165 (1966); and M. Shimoji and H. Hashino, J. Phys.

Chem. SoHds 28, 1155, 1169 (1967).

[65] F. Liity, in The Physics of Color Centers, ed. W. D. Compton.

to be published.

[66] R. Tucker, A. Laskar, and R. Thomson, J. Appl. Phys. 34.

445 (1963); see also various recent papers of Ya. Geguzir and co-workers, such as Sov. Phys. — Dok. 11, 626 (1967): Sov. Phys. -Solid State 8, 2599 (1967); ibid. 7 , 282e

(1966).

[67] P. SUptitiz and J. Teltow, Phys. Stat. Solidi 23, 9 (1967)

7

Diffusion of Sodium in Sodium Chloride in an Applied Electric Field ^

Robert J. Friauf and Vaughn C. Nelson

University of Kansas 66045

1. Statement of the Problem

Ionic conductivity and diffusion in NaCl are both

due to the movement of vacancies. Conductivity

and diffusion can be compared by the Einstein

relation, corrected for correlation, but there are differences which can be attributed to diffusion by

vacancy pairs. Vacancy pairs are electrically neutral

and would not contribute to the conductivity. If the drift mobility of the cation could be measured

by diffusing Na in an electric field, the amount of

vacancy pair diffusion could be obtained from this measurement and checked against the measured

conductivity and the Na and CI diffusion without an electric field.

2. Procedure

(a) The ionic conductivity was measured in the

extrinsic and intrinsic ranges (275 to 797 °C) with an ac bridge at 1000 Hz. (6) The diffusion of Na was measured (640 to 790 °C) with radioactive tracers and the sectioning method, (c) The mobihty of

radioactive tracers was measured (600 to 772 °C)

by diffusing Na in an applied electric field. The

amount of charge passed through the sample was

measured with a silver coulometer. Considerable

difficulties were encountered because of the poor

contact at the interface between the two crystals.

' Work supported by the U.S. Atomic Energy Commission.

^ Present address: Universidad de Oriente, Cumana, Venezuela.

3. Results

(a) The Na diffusion (no electric field) is given by

£>= 1300 (cmVsec). (b) The diffusion of Na

in an applied field gives a drift mobility of Na ions

Mp=(lXlO«/r) e-2»6/^-^ (cm'Vvolt sec), (c) The

best two exponential fit of the conductivity is

(tT= 3.6 X 10« e-2 «6M-^+ 2.5 X 10» e-2-08Mr

{°K ohm"' cm"')

where the cation parameters in the first term have

been calculated from the measured drift mobility

of Na tracers, (d) The activation energy for diffusion of Na by vacancy pairs is 2.6 eV, and approximately

50 percent of the Na diffusion near the melting

point is due to vacancy pairs.

4. Conclusions

Measuring mobihty in alkali halides by diffusion

in an electric field is promising as a tool in actually

delineating the contributions of single vacancies.

The importance of anion vacancy conductivity in NaCl is confirmed by these results. The overall picture for NaCl shows an appreciable contribution

of vacancy pairs to the diffusion of both anion and

cation.

290-621 0-68—

2

9

Characterization of Point Defects in Oxides ^

John E. Wertz

Department of Chemistry, University of Minnesota, Minneapolis 55455

This review describes the geometry of numerous point defects which have been observed in oxides. Experimental evidence for their existence and for their configuration is briefly cited. Attention

is concentrated primarily upon the alkaline earth oxides, for which most defect structural information

is available. Some defects in ZnO and in AI2O3 are also mentioned. Defects may be categorized as involving impurities, vacancies, or interstitial atoms. These may exist in isolation or in association with other defects. Numerous examples are given of cation vacancies associated with one or with two

other defects; the latter are neutral in each case. One may further categorize defects as to whether

they require radiation of the solid for their production. Valuable information about some stable defects has been obtained by correlating them with closely related unstable defects. Some defects may be

studied by optical or electron spin resonance techniques if they are caused to trap electrons or holes. An example is the determination of the site geometry of OH~ ions which in MgO show an infrared band at 3296 cm Measurement of the concentration of trapped-electron-or trapped-hole centers after appropriate sample treatments may give minimum estimates of the numbers of cation or anion va- cancies present. For some defects, electron spin resonance studies have permitted correlation with

optical absorption, luminescence or uniaxial stress measurements. Known trapped-hole and trappedelectron defects in oxides are enumerated.

Key Words: Electron spin resonance, impurities, interstitials, optical spectra, point defects, va- cancies.

1. Introduction

The perceptive experimentalist or theorist who

concerns himself with mass transport in oxides

recognizes the significance both of hne- and of

point-defects. Notwithstanding his recognition of

the latter, he may disdain confrontation with defects resulting from various kinds of irradiation.

If nature had provided us with observation capabilities of atomic resolution for direct detection

of such defects as vacancies, one might not have

to resort to irradiation and other procedures for their characterization. It is a simple fact that the

geometry of a number of stable defects in oxides

is known because they were converted by irradia- tion to unstable defects amenable to study by

optical or electron spin resonance methods. In

the process, it has been possible in some cases to provide unambiguous analytical procedures for measuring the numbers of the stable defect. Alternatively, for some oxides these studies make

it clear that it is naive to refer to Schottky pairs,

for anion vacancies have been shown to be far outnumbered by cation vacancies at temperatures

below 1200 °C. To demonstrate this, it was necessary to create anion vacancies by neutron irradiation, populate them singly with electrons, and

' Some previously unpublished results cited herein were obtained with the support

of the Solid State Sciences Division of the Air Force Office of Scientific Research

under grant 200-66.

identify them with one or more measurable properties. Thus it was estabhshed that not even

ionizing radiation could generate such vacancies

at room temperatures (at least in the alkaline earth

oxides). In the conviction that many students of mass transport will wish to know which defects

may be detected in oxides as well as the nature of

the evidence for their existence, this review is presented without apology for inclusion of unstable

point defects. Outside of the group of alkaline earth oxides,

there are relatively few oxides in which the geometry

of point defects (except perhaps impurity ions) is known with reasonable certainty. The simple rocksalt structure and the ready availability of single

crystals of MgO, CaO, and SrO makes these —

especially the stable MgO -ideally suited for characterization of defects. In a number of instances,

the characterization of a particular defect in MgO

led to reports of observation of the analogous defect

for another type of oxide. Hence much of this review is devoted to defects in MgO.

2. Substitutional and Interstitial

Defects

2.1. Substitutional Impurities

The nature of the impurities present in a sample,

their distribution and their interaction with other

11
defects are vital to one who seeks to make meaningful measurements of mass transport in oxides. In

ionic crystals, the impurities are predominantly

found in substitutional sites. Even if the charge on an impurity ion differs from that in the host lattice, one may still find a significant fraction of ions

without a compensating impurity or vacancy nearby.

Such statements may be made with considerable

confidence if the ion belongs to transition metal or

rare earth groups, in a paramagnetic state readily

studied by electron spin resonance. For such ions

the number and anisotropy of the lines in the elec- tron spin resonance (ESR) spectrum reflect the

electric field symmetry about the ion. If the angular

dependence of the spectrum is precisely that which

is expected from the symmetry of the host, the ion

is likely to have no nearby compensating defects. The existence of ions without an associated compensating vacancy is especially common in the

oxides of formula MO. Here one cation vacancy

compensates for two trivalent ions, and at least one of these may be remote from the vacancy.

Monovalent cations may serve to compensate impurity cation charge in some oxides. Even if one

detects no departure from the octahedral or tetra- hedral symmetry of the host ions, there may be a symmetrical distortion of the lattice about an impurity ion because of a different fractional covalency

of an impurity-oxygen bond.

The ESR behavior of ions with an associated

compensating defect is considered in section 3. It is important to recognize that a number of the

transition metal ions may exist in two or more

valence states. An outstanding example is iron,

which in the alkaline earth oxides may exist stably

in the divalent or trivalent forms; irradiation leads

to a significant concentration of monovalent iron. No attempt will be made here to enumerate the

substitutional transition-metal or rare-earth impurity

ions which have been identified in various oxides

by ESR techniques. A compilation has recently

been made of ESR parameters for ions in MgO,

CaO, SrO, AI2O3, TiOa, and ZnO, as well as some compounds having perovskite, spinel, or garnet

structures [1].^ However, section 3 is devoted to those impurity ions for which there is evidence of

association with vacancies or with other defects,

both stable and unstable.

Very little is known about the sites occupied by

a number of impurities detected by standard analyti- cal techniques. Many of these are diamagnetic and

further have not modified the ESR spectrum of a paramagnetic defect in so obvious a way as to an- nounce their presence. They are especiaUy likely

Figures in brackets indicate the literature references at the end of this paper.

to remain undetected by optical methods as well

as by ESR techniques if there is no abundant

nuclide of nonzero spin. One long-known defect

«;hows the presence of aluminum in association

with a vacancy at which a hole is trapped by its hyperfine splitting pattern [2]. More recently, the

effects of an aluminum ion associated with a vacancy were detected by its reduction of the zerofield spUtting of a Cr^+ ion also associated with

the vacancy [3, 4, 5]. This defect is considered in section 3. The association of a lithium ion with a cation vacancy in ZnO is obvious from a hyperfine

quartet from ^Li (1 = 3/2) which becomes a triplet

from ^Li when the sample is doped with the latter

[6, 7].

2.2. Interstitial Defects

Relatively few interstitial defects have been

established with reasonable certainty in the oxides.

The presence of interstitial zinc in ZnO heated at elevated temperatures has long been accepted.

In the alkaline earth oxides one infers the presence

of interstitial alkaUne earth metal atoms and of oxygen atoms only after neutron- or ion-irradiation. However, there seems to be Uttle direct or detailed

evidence of such atoms. For one impurity, the 0H~

ion, remote from a cation vacancy, it seems rea- sonable to infer an interstitial position for the

hydrogen atom of the OH" ion [8]. This location is consistent with the relatively great width of the

infrared band (~ 100 cm"^) presumed to be cor- related with it; the widths of OH stretching bands

for an O • • • H • • • O system increases with reduced

separation of the oxygen atoms [9]; for hydrogen

between two oxgyen atoms at approximately the

normal lattice spacing of 4.20 A, the width of the

infrared band is roughly one-tenth as great. These

defects are described in section 4.

3. Association of Vacancies With

Impurities

The defects of greatest interest for mass trans- port studies in most oxides are the anion and cation

vacancies. As pointed out earlier, it is erroneous

to assume that these are present in equal concentrations in oxides except perhaps at very high temperatures. One presently has no technique for direct observations on isolated vacancies in oxides.

However, an estimate of their number may be made

if the cation vacancy is caused to trap a positive

hole or an anion vacancy an electron. These types

of defects will be described in sections 4 and 5

respectively. Evidence from trapped electron

12
defects will also be adduced for the presence of a divacancy and of a trivacancy (section 5). Vacancies which are associated with other

point defects may be illustrated for the alkaline

earth oxides as follows:

a. Cation vacancy-trivalent cation, impurity as next-nearest neighbor (nnn).

b. Cation vacancy-trivalent cation, impurity as nearest neighbor (nn).

c. Cation vacancy-monovalent anion impurity.

d. Cation vacancy plus two trivalent cations,

each as next-nearest neighbor.

e. Cation vacancy plus two trivalent cations,

each as nearest neighbor.

f. Cation vacancy plus trivalent cation plus

monovalent anion.

g. Cation vacancy plus trivalent cation plus

trapped hole.

h. Cation vacancy plus monovalent anion plus

trapped hole. These various defects are illustrated in figure 1 for an alkahne earth oxide. For simphcity, only

deviations from the normal site charge of the

rocksalt-type lattica are shown. Thus, M+ repre- sents a trivalent impurity cation, A+ a monovalent anion and 0H+ and 0+ represent the ions

OH" and O" respectively. The symbols [=J and

refer respectively to unoccupied cation and

anion vacancies. Since each of the associated

defects considered here is a hnear array of simple

defects, only the minimum number of atoms will henceforth be used to represent them in the text. The relative proportions of isolated, singlyassociated and of doubly-associated vacancies ob- served at room temperature depends upon the

temperature at which equilibrium is attained if quenching occurs at a sufficiently rapid rate. In MgO, there is an appreciable mobihty of cation

vacancies at 400 °C, and hence if isolated defects

are to predominate, one should rapidly quench

crystals from some temperature of the order of

1000 °C. In the temperature region near 400 °C,

a considerable association of vacancies with one

or more impurity ions will occur [4]. One may

obtain binding energies of associated defects by

studying their concentrations at room temperature,

following a quench from successively higher

temperatures [10]. Anion vacancies are normally

present at very low concentrations in the alkaline

earth oxides. However, if they are produced by

neutron irradiation in a reactor, one may follow the

disappearance of isolated vacancies and detect

the presence of increasing numbers of vacancy

pairs as the temperature is increased up to 900 °C. This association is detected from the disappear- ance of F-centers and the formation oi F 2 centers

(sec. .5).

3.1. Structural Data from Electron Spin

Resonance

The centers

Cr+ O [=] [001] or equivalent

and Cr+ [=] [OH] or equivalent

in MgO were among the early point defects in oxides

for which there was evidence of association of a vacancy with another defect [11, 12]. As in figure 1, only deviations from normal site charge are shown.

The first defect, with a vacancy in the nnn position

has a principal axis of the crystal as an axis of

tetragonal symmetry. For brevity, this will be

referred to as a Cr? defect. The second defect, with

a nn vacancy, shows an ESR spectrum interpretable

in terms of rhombic symmetry; it wiU be called a CrR defect. Additional there are Cr^+ ions which

show an isotropic ESR spectrum, and are therefore

OoOoO;oOoOoO

3 oOoOo!o®'OoOo e O ®*OHO!o OSO o O

_ o_0 o_o_ojO o O ®10_o_

OoOoOloOoOoO

o O ®*0 op®*OH© o

OoOBOioOoQoO

_ <LCIo O o !o oQo O o_

OoOoQIoOoOoO'

c o OS® ojO ®"OH® o g

_0_o O o_0]o 0 o_Oj) 0

oOoo^oioo^o^o'"

0®'OBO®10oOoO

d oOoOoO|0@B®o h

OoOoOojOoOoO oOoOoQoOoOo

Figure 1. Geometries of known associated cation-vacancy point

defects in alkaline earth oxides. One cation vacancy is associated respectively with:

a. One nnn trivalent cation, (001) or equivalent axis.

b. One nn trivalent cation, [Ollj-type direction.

c. One nn monovalent anion.

d. Two trivalent nnn cations.

e. Two trivalent nn cations.

f. One nnn trivalent cation and one nn monovalent anion.

g. One nnn cation and one nn positive hole.

h. One nn anion and one nn positive hole.

13
regarded as unassociated ions in octahedral symmetry. These wiU be represented byCro.

The ESR spectrum of the Cro ion consists of a

single line with g= 1.980 for the even isotopes of Cr. A hyperfine quartet from the ^^Cr in its natural

abundance of 10.53 percent shows a characteristic

splitting constant of about 17.9 G [11]. Upon association with a nnn cation vacancy in MgO, the A/s = ±l/2 states of Ct^+ are split from

the Ms — ±312 states by about 0.16 cm~* in zero magnetic field, due to the lowering of the symmetry

from octahedral to tetragonal. This splitting is given by 2D in the spin HamUtonian for Cr^+

(S = 3/2):

<r = /3 [g'x//xSx + gyHySy + g,H,S,]

+ /)[S|-S(S+l)/3]+£(S2-S2). (1)

Here gx = gy^ g^-, gz^ g\\, and E = ^ because of the

tetragonal axis of symmetry. The g components for the several ions associated with vacancies are given in table 1, together with zero-field splitting

parameters D. The ESR spectrum is shown in figure 2 with the Hues belonging to the Crr center

labeled with a T. For Cr^^ with a nn cation vacancy, the site symmetry is rhombic. Hence the j^'-components will not be identical, and the asymmetry parameter E will be different from zero. The zero-field splitting

resulting from the close proximity of the vacancy

to the Cr'+ ion in MgO is about 0.8 cm''^ Since

the microwave photon energy for the spectrum

shown in figure 2 is only about 0.3 cm"\ one ob- serves only a few of the allowed transitions occurring

where energy levels come sufficiently close together.

These are marked with an R in figure 2. If the Cr^+ concentration is small, i.e., < 0.01 percent, one expects that most of these ions wiU

be widely separated in samples quenched from

perhaps 1000 °C. Because of the double charge of

the vacancy, one must serve for two Cr^+ ions. Hence for low concentrations, only half of the

Cr^"*" ions may be associated with a vacancy. The

normal relative populations are:

[Cro]>[Crr]>[Cr«],

where the square brackets imply concentration of

the enclosed species. Only recently has the ESR

spectrum of the nnn Crr center been observed in CaO [13]. The rhombic Ctr defect has not been

reported in any crystal other than MgO.

Analogous associated vacancy centers have been

observed for Fe^+ in MgO, although ordinarily the

concentrations of nnn or nn centers is at least an

order of magnitude smaller than that of Fe^+ in octahedral symmetry [14]. It is not clear why the

relative fraction of Fe^+ associated with vacancies

is less than that for Cr^+. The ESR parameters are given in table 1. Association of vacancies with Ti^+ represents an

interesting case [15]. Unlike iron, which may be

incorporated in either the + 2 or + 3 states, both of

which are detectable by ESR spectra, Ti appears

to be incorporated in the tetravalent form in MgO.

One therefore requires x-irradiation to create

Ti^+ {"id}) ions. If any of the latter were simply

substitutional ions in octahedral symmetry, they

would not be detected in the ESR spectrum because

Table 1. ESR parameters for associated vacancy centers in MgO

Geometry g\\ g±

|D|

(X 10^ cm"'

|E|

)

Ref.

Cr- 0 [=] [001]... 1.9782 819.2 [11] CrM=] [Oil]... 338.5 2511 [12] Fe+ 0 [=] [001]... 2.0037 5000 [14]

Ti+0 [=] [001]... 1.9533 1.8993 [15] Mn++ 0 Li-? [001]... 1.995 280 [5] Mn++ 0 [=] [001]... 1.9926 528.3 [13]

Cr+ 0 [=] 0 M+ [001]... 1.9782 798.3 [4, 5] Cr+ [=] M+ [Oil]... 331.1 2207 [5]

Cr+ [=] H+0 [001]... 1.9786 1.9815 746.5 [13]

For CaO:

Cr- 0 [=] [001]... 1.9697 1.9751 1360.6 [13]

14
T4 Tg

I O

I

Figure 2. £5^ spectrum of Ct^+ in MgO.

Lines marked with symbols O, T, R, T' and R' refer respectively to Cr"* in octahedral symmetry, to defects a, b, d and e of figure 1. The unusual aspect of this spectrum is the

inversion of relative intensity of primed and unprimed defects. This results from heating at 400 °C to enhance vacancy association.

of the orbital degeneracy of the ground state. However, the tetragonal crystal field distortion from the nnn vacancy lifts the degeneracy. It seems reasonable to expect that unless samples were quenched

from a high temperature (perhaps > 1200 °C) that most Ti^^ ions should have an associated vacancy.

The ESR spectrum of Mn*+ in a site of tetragonal

symmetry has been observed for Li-doped MgO

powder; the parameters are given in table 1 [5].

However, the zero-field sphtting parameter D is very much lower than that of any of the other ions.

It is suggested that perhaps in this case Li+ is the

associate of Mn'*+ [5] . Perhaps one may write

Mn2+ O Lias the configuration. Another Mn^+ spectrum of

tetragonal symmetry has been observed in a single

crystal [13]. The magnitude of the parameter D

(table 1) makes it appear that this is hkely cor- responds to the defect

Mn2+ O [=].

In addition to the Ctt or Cr^ centers which have

a net negative charge because of overcompensation by the associated cation vacancy, there are

the neutral, doubly-associated centers shown in

figure Id and le. Evidence for these in the ESR

spectrum is found in the form of satellites to the

Ctt or Cth spectra of figure 2. These satellite spectra contain precisely the same number of lines as the Crr or Ctr spectra respectively, but with

the slightly smaller zero-field-spUtting constants

given in table 1 [3, 4, 5]. These are interpreted as indicating a reduced crystal field at the Cr^+ ion

because of another trivalent ion symmetrically

located on the opposite side of the vacancy. The

Unear arrays for these centers are thus:

Crr-: Cr+OHOM^

Cr«': Cr+[=]M+

( 001 ) axis [011]-type direction.

In an aluminum-doped MgO crystal the ion M+

is very hkely to be AF+. By extended heating of

this crystal at 400 °C, it was possible to obtain a Crr spectrum 90-fold as intense as the Crr spec- trum. This temperature of 400 °C has previously

been mentioned as lying within the range in which

there is a significant mobility of positive ion va- cancies. Subsequent heating at 700 °C reduces

the Crr intensity many-fold. The Cr^- concentration changes qualitatively in the same way as

that of Crr. In undoped crystals it was possible to get a ratio of Crr/Crr concentrations of about 2.2.

15
In MgO crystals which are known to have a

significant concentration of the defect

0[=]H+0

which shows an OH-stretching line at 3296 cm~'

[16, 17], one sees still a third Cr*+ spectrum of

tetragonal symmetry. It appears to be the neutral

defect

Cr+OHH+0.

Here the compensating charge is still closer to the vacancy than in the Crr center and the zerofield splitting constant (table 1) is the smallest of

the three doubly-associated centers involving

Cr3+ [13]. In the infrared spectrum of MgO, the OH Hne

at 3296 cm"' is always accompanied by another

at 3310 cm"' [16, 17]. The latter hne appears to increase with the concentration of trivalent ions

in the crystal. If any such trivalent ion is asso- ciated to form the center M+0[=]HO+,

then one may understand the lack of a change in intensity of the 3310 cm"' line on ;c -irradiation of

the sample, since the center is neutral [8]. This

behavior is unUke that of the defect

0[=]H0+

which may trap a positive hole on the oxygen atom

opposite to the H0+, and thus shift the OH stretch- ing frequency from 3296 cm"' to 3323 cm"'. It re- mains to be estabhshed that there is a corre- spondence between the Cr+0[=]H+0 center, for which the ESR spectral parameters are given in table 1, and the one responsible for the 3310 cm"'

Line. Presumably AF+ might be the trivalent ion

in some of these defects.

It is known that there may be an appreciable concentration of F" (fluoride) ions in some crystals.

Presumably they may also associate themselves

with a cation vacancy, for one may detect a defect

with the structure O"*" [=] F+ (sec. 4).

3.2. Associated Vacancy Data From

Luminescence

Some valuable structural data on point defects

of Cr^+ may be obtained from sharp-line lumines- cence studies on oxides. A pair of very narrow fines

("/? fines") observed in ruby (Cr'+'doped AI2O3)

corresponds to ^Ea ^A^g transitions. In the alkafine

earth oxides, a single fine at 6981 A is observed

for octahedral symmetry [18—21]. (Phonon sidebands are also present, but they wiU be ignored

in this discussion.) It has been estabfished that in MgO the fine from Cr^+ ions in octahedral symmetry may be spfit by the action of a uniaxial stress

[22]. If the symmetry about the Cr^+ ion is lowered

to tetragonal, both the ^Eg and *Aig states are

spfit. The magnitude of the ground state spfittings

is given from ESR data (in the previous section)

by 2Z> = 0.16 cm"' [11]. This spfitting is so smaU

as to be unresolved even at 4 °K; it serves to broaden fines relative to tfie R fine. At low concentrations of chromium (~ 10"^) in MgO there is a weaker pair of fines at 6992 A and at 7038 A; their

intensity increases with Cr^+ concentration. These

are attributed to Cr^+ ions associated with a nnn

cation vacancy, i.e., the Crr center [20]. Uniaxial

stress experiments verify the tetragonal symmetry

[21]. External stress along the axis of the defect causes additional spfitting of the two fines, and

hence the internal stress is compressive. That is, the vacancy-Cr^+ distance is less than the normal

lattice constant of 4.20 A. No spfitting occurs for a stress appfied along a [lll]-direction. The ratio of the intensities of the fines at 6992 A and 7038 A

varies with temperature according to the Boltzmann

factor. This is strong evidence that the 6992 and

7038 A fines arise from a 94 cm~' spfitting of the % state [21]. A pair of emission fines occurring at 6989 and

7034 A are very weak at low Cr^+ concentrations

but their intensity increases rapidly with concentration. Uniaxial stress experiments show that the

symmetry axis is a principal axis of the crystal [21]. Likewise, no spfitting is observed for stresses appfied along a [lll]-direction; relative intensities

of 6989 and 7034 A fines again vary with temperature

approximately as the Boltzmann factor. These fines have been attributed to a doubly associated center

involving two Cr^+ ions, i.e., Cr+ O [=] O Cr+. A

smaU exchange interaction would be expected for two Cr^+ ions thus disposed, and structure about

both the 6989 and 7034 A fines has been interpreted

as evidence for such interaction. Comparison at 18 °K and at 4 °K of the intensities of the satelfites about the 7034 A fine suggests a 2.0 cm~' exchange

spfitting of the "^Eg level and a 2.6 cm"' spfitting

of the *^2g ground state [21]. An additional Cr^+ luminescence fine has been

reported at 7088 A in MgO [10]. Its relative lumines- cence intensity is reported to be correlated with

that of the ESR spectrum of the Qvr center. The

large shift in position from the 6981 A fine for Cr^+

in octahedral symmetry is quafitatively to be ex- pected from the large distortion of the crystal field by the nn vacancy. Hence one appears to have lumi- nescence data which are in accord with expectation

for defects la, lb. Id, and le of figure 1.

16
Ions which are isoelectronic with Cr^+ might be

expected to show characteristic -Eg—**A2g hnes

appropriate to the octahedral symmetry of an

isolated substitutional ion or a distortion to lower

symmetry by vacancy association. The emission

line of in octahedral symmetry in MgO at 8696 A at 77 °K is accompanied by a group of phonon sidebands; it corresponds to the Cr^+ Hne

system originating at 6981 A [23]. No emission hnes

corresponding to a V'-'^ ion associated with a cation

vacancy has yet been estabhshed. x-irradiation

does result in electron trapping by numerous

V^^ ions in octahedral symmetry. It appears

hkely that luminescence hnes reflecting va- cancy association resulting from the parent

ion will yet be established.

The Mn*+ ion fluorescence or cathodolumines- cence spectrum in MgO shows numerous hnes

[5, 8. 17]. Of these, the hne at 6545 A is attributed

to ions in octahedral symmetry [5]: another at 6593 A is attributed [8] to the center Mn^+OHIt remains to be estabhshed by correlation with

the ESR spectrum of Mn^^ in tetragonal symmetry

that the latter assignment is correct.

4. Trapped-Hole Defects

The term "color center" apphed to the alkali hahdes is hnked with defects having an electron

in an anion vacancy, i.e., "F-centers." Such defects may readily be generated by irradiation or by heating in metal vapors. Though comphcated

by the relatively more pronounced effects of impurities, the alkahne earth oxides do show a characteristic color center on irradiation. This center

is the hole trapped on an oxygen atom adjacent

to a cation vacancy (fig. 3) [24]; it is here referred

to as the Fj center. Though the model was proposed

for the alkah hahdes [25], it has not been observed

in them.

At 77 °K, MgO crystals which have been subjected to 2537 A or x-irradiation show an ESR

spectrum of three hnes at random orientation of

the magnetic field [24]. The angular dependence

of these hnes is that for a system of tetragonal

symmetry. The positions of the Hnes at several

field orientations is given in figure 4. The ESR

parameters for this defect and for others akin to

it are given in table 2. The positive deviation of the ^'-components from the free-electron value and

the tolerance of high levels of microwave power

are consistent with the assumption of a positive

hole located on an oxygen atom adjacent to a cation

vacancy, i.e., an O" ion. From the width of the

Oo GoG o

o © BGo ©

Go GoG o

o G o

©&b

o'

Go o

oG oGo o_ Go O o O o

o © B©o G

Go GoG o

o"G' oQo o~

Go ©@© o

oG o © o G

Models of trapped holes in alkaline earth oxides.

a. Vi center.

b. Vqh center.

c. Vp center.

d. W center.

V = 9.065 GHz

I I H II [001] 3173 32 30

I Hllloii]

3173 3201

III]

H

Figure 4. Line positions in the ESR spectrum of the Vi center

in MgO for H||[001], H||[011] and H||[lll]. The variation of line position as a function of angle between the. magnetic field and

a principal axis is shown at the bottom of the figure.

hnes, which may be as httle as 0.4 G, one may infer that the lifetime of a positive hole along a particular axis is long compared with 10~' sec. This

relatively long hfetime is presumably the result of

17
Table 2. ESR parameters of Y-centers

Vu MgO

F,, CaO

f^oH, MgO

Vp, MgO

W, MgO

ZnO (Li-doped).... AI2O3

Al203(0+?)

2.0032

2.0011 2.0033

2.0031 2.0032 2.0026

2.02 2.018

g±

2.0385 2.0710

2.0396

2.0388

2.0408 2.0254

2.006

2.011

D

(cm-')

0.02125

Ref.

[24]

[27]

[16, 13] [28]

[24, 13]

[7] [36]

[29]

an extensive relaxation along the hole-vacancy

axis. At room temperature the lines are roughly

thirty-fold broader, suggesting a much more rapid

hole migration among the three principal axes. While this representation of the Vi center geometry

in MgO gives the impression of localization of the

hole upon one oxygen atom, the Fi-ESR data re- quire only that the fractional localization of the

hole on the two axial atoms differ from that on the

four equatorial atoms.

If the lifetime of a hole upon a particular oxygen

atom v\rere at least of the order of seconds, one

might hope selectively to populate or bleach Vi centers along a particular axis. Failure of such

experiments even at 4 °K suggests that there is considerable hopping of a hole about a particular

cation vacancy [13, 26].

That MgO becomes violet colored upon irradia- tion was reported in 1937 [30]. The color corre- sponds to a broad absorption band at 5320 A. Since

the coloration of several specimens varies widely

for a given irradiation, it was concluded understandably that the coloration is due to an impurity.

However, parallel optical absorption and ESR ex- periments have shown that the 18000 cm ' band is that of the center [31]. In some samples the

crystals appear brown in varying degrees after

jc-irradiation. Such behavior is attributable to an

additional absorption band appearing at 3940 A. This band may be induced in any MgO crystal by

heating in air at 1200 °C. Its origin is unknown.

While the negative ion vacancy of the F-center may be produced in the alkali halides by ionizing

radiation, neither a cation nor an anion vacancy is generated by such irradiation of the alkahne earth

oxides. However, positive ion vacancies occur in significant numbers in all magnesium oxide crystals which have been available to this time; these

are the inevitable counterparts of trivalent impurity ions which are always present. The aluminum

ion is an important contributor of cation vacancies,

though it does not make its presence as obvious

as the paramagnetic ions Cr'^ and Fe^+. The two

latter ions (and others derived from them on irradiation) affect a considerable number of the

physical properties of MgO. Other impurities in MgO which may help to incorporate positive ion

vacancies are Ti'*+, and Mn'*^. Iron is doubtless

the most important impurity controlling the number of positive ion vacancies. In MgO crystals

grown from the melt, the predominant fraction of

iron is in the Fe^"*^ form, as judged both by the

ESR- and the optical absorption spectra. Extended

heating in oxygen at 1200 °C converts a significant

fraction to Fe^+ [32], presumably accompanied

by the inward diffusion of cation vacancies [33]. We have noted that the cation vacancy may be

associated with defects such as Ti'^^, Cr^+, Fe^+,

OH , or F". Heating to 700 °C or above and quenching may serve to dissociate such centers and

increase the number of isolated positive ion va- cancies. Hence such heating and quenching before

irradiation with x rays or with ultraviolet light of

energy 5 eV or more gives an enhanced concentration of Vi centers. It has tacitly been assumed

that there exists also a number of defects adequate to accept the electrons removed by UV

or x-irradiation [26]. Probably most of the holes

are generated at normal oxygen sites. These would move rapidly through the crystal until trapped by

a transition metal impurity ion or by a cation

vacancy.

Before proceeding to a description of the ESR

spectra of other hole centers for which the spectral

parameters are given in table 2 and the models are given in figure 3, it is profitable to consider data

concerned with the distribution of hydrogen as an

impurity in MgO.

A narrow 0-H stretching band appears at 3296

cm"' in many MgO crystals before these have undergone any treatment. It is accompanied by a band at 3310 cm^' only partially resolved from it and ap- pearing as a shoulder [16]. Other crystals merely

give a very broad band in this region upon heating

at 400 °C. These show an apparent transformation

of the broad band into the 3296-3310 cm"' pair,

in addition to weaker fines. In some cases one sees

a strong fine at 3700 cm"'. The close correspondence

of this fine with that in Ca(0H)2 suggests that one may have local regions of Mg(0H>2 in an MgO

crystal. This has been confirmed from the dif- fraction pattern in MgO of a small region observed

to diminish in size under the electron beam of an

electron microscope [34]. The broad band in the

OH-stretching region presumably arises from OH"

in isolated sites. For these the hydrogen atom

would probably be interstitially oriented [8]. These

infrared bands are especially strong in crystals

which have been made from MgO powders pre- pared from Mg(0H)2. Crystals showing strong

3296 cm"' bands are also found to produce lentic18
ular cavities when heated at 1000 °C in a reducing

gas such as H2, D2, or CO. The source of the

hydrogen appears to be primarily derived from OH

: 1 centers having the 3296 cm"' hne [34]. Upon x-irradiation of such samples at 77 °K one

observes a reduction of intensity of the 3296 cm~*

i line by a factor of two or more, while a new line

! appears at 3323 cm~'. The latter has an intensity

I approximately equal to that lost by the 3296 cm~'

\ line. At room temperature the 3323 cm"' hne

bleaches with a time constant of 80 min at 32 °C,

j with a corresponding regrowth of the 3296 cm"'

line. The interpretation of the infrared hne be- havior is based upon a correlation with ESR

behavior.

In all crystals showing a strong 3296 cm"' line, one sees doublet satellites on one side of each of

the Vi center ESR hues. These doublets decay

with the same time constant as that of the 3323

cm"' infrared line. Their angular dependence

hkewise indicates a center of tetragonal electric

field symmetry. The suggestion is strong that the

defect is a hole trapped at a positive ion vacancy,

but with a perturbing nucleus of spin one-half. This nucleus is required by the axial dependence

I of the hyperfine sphtting to he along the axis deI fined by the positive hole and the cation vacancy.

Electron-nuclear double resonance experiments

unambiguously demonstrate the nucleus to be

hydrogen. The magnitude of the anisotropic component of the hyperfine sphtting is a measure of

the dipole-dipole interaction of the positive hole

and the nucleus. From this interaction one may

calculate r=3.2 A, confirming the proton position

indicated for the FoH-defect in figure 3, i.e., adjacent

to the cation vacancy. This distance is fortuitously close to that obtained by assuming that the oxygen

atoms are undisplaced from their normal distance

of 4.20 A and that the OH distance is 0.94 A as

in Ca(0H)2. The intensity of the ESR hnes of this center correlates well with that of the 3323

1 cm"' hne. Decay of the Fqh center is undoubtedly

i due to the loss of the hole from the oxygen adja- cent to the cation vacancy. Hence the transfer of

I line intensity from the 3323 cm"' to the 3296 cm"'

I hne of decay of the Vq^ center implies that the

' 3296 cm"' hne is appropriate to the geometry

0[=]H+0. Here, although the Fqh center pro- duced by irradiation was the primary defect of

ESR study, the by-product — more important

I perhaps than the Fqh center — was the unambig-

' uous identification of a stable defect present in

1 significant concentrations in crystals subjected to

I a wide variety of studies by many investigators.

Even in some MgO crystals which have not been

x-irradiated there may exist weak 3323 cm"' hnes

' which suggest the presence of Vqh centers [8, 13].

Their presence in Li-doped MgO was interpreted

as due to raising of the Fermi level by lithium [8].

It has been surmised that the 3310 cm ' infrared

line arises from OH" adjacent to a cation vacancy

and associated with a trivalent impurity cation,

corresponding to the model of figure If [8], i.e., M+0[=]H+0.

The ESR data for a center in which the trivalent cation is Cr'^^ were given in section 4, and AF"*" was mentioned as another possible cation. Such

an assignment is consistent with the observation

that the intensity of the 3310 cm ' hne is unaltered

upon x-irradiation [16]. Another defect which may be present in nonirradiated MgO has an analogous structure and

could be described in detail because a trapped

hole could be associated with it. This trapped hole

defect is labeled Vp in figure 3, and has the configuration [28]

0+ [=] F+

Again, a doublet, FoH-hke ESR spectrum is ob- served. The identification was made when the

g'-components of HF-treated MgO powder samples were found to coincide with those of the spectrum

in a single crystal. Fluoride ion is a common impurity in MgO, and within the range in which the

temperature is high enough to permit cation va- cancy migration but low enough to permit stable association, one may expect considerable association of the two defects. Following estabhshment of the Vi center in MgO,

it is obvious to ask about the possible occurrence

of a center formed by trapping two holes at a single

cation vacancy. This W center (i.e., double-V)

shown in figure 3 has indeed been observed in MgO x-irradiated and observed at 77 °K [24]. The

ESR spectrum of this system with S = 1 may roughly

be described by the statement that each hne of the Vi center is replaced by a pair of widely sep- arated hnes. Having a spin of 1, the energy levels of the system are spht in zero magnetic field, with

D = — 0.02125 cm"'. From the magnitude of the

dipolar interaction one calculates a separation of 4.9 A for the two holes. The axis is again along a

[001] or equivalent direction. It is apparent that on

a locahzed model the two holes tend to be as far apart as possible. This center is far less stable than the Fi center, probably due to a reduced binding energy for the second hole. It is the least stable of the various neutral centers which contain a

cation vacancy.

There appears to be one other example of a pair of holes trapped on two oxygen atoms adjacent to

19
a cation vacancy. This is observed in a reactor- irradiated AI2O3 crystal in w^hich at 77 °K for Hl{c), one observes a pair of lines for each of

six inequivalent orientations of a line joining two

oxygen sites [35]. For H\\{c), only two lines are observed either at 77 °K or at 300 °K. However,

for H^{c), as one raises the temperature above

230 °K, the twelve lines are observed to broaden

at an increasing rate. This is attributed to jumping

of the holes among the equivalent positions. For

this center with S=l, Z) ~ 0.13 cm ^ From this D value, an internuclear distance of about 3.0 A

has been computed for the two interacting holes,

in satisfying accord with expectation. The same

spectrum had previously been given a more doubtful interpretation [36].

It is reported that when hthium is substituted

for zinc in ZnO that a positive hole may be trapped on an oxygen atom adjacent to the hthium [6, 7]. The hyperfine sphtting pattern is appropriate to the overlap of the hole wave function on the hthium.

Doping with ''Li gives rise to the expected triplet

sphtting for *^Li with 1=1 [7].

5. Electrons Trapped at Anion

Vacancies

5.1. Electron Spin Resonance Evidence for

F-Centers

The first documented example of an electron

trapped in an anion vacancy in an oxide appears

to be that in MgO [37]. Subsequently, ESR ab- sorption of similar centers was detected in CaO,

SrO and BaO [38, 39], ZnO [7], AI2O3 [36] and in some complex oxides [35, 40]. ESR data for electron centers in some of these oxides is given in table 3. The ESR spectrum in MgO has an isotropic g- factor of 2.0023. Since most of the Mg^+ neighbors

of an anion vacancy are -^Mg or '^^Mg with 1=0,

there is an intense central hne. However, for a

natural abundance of "Mg of 10.11 percent, the ] probabihty is about 0.36 that one of six Mg neighbors is -^Mg with 1 = 5/2. Hence one would expect

a hyperfine sextet, but with marked anisotropy.

Since the axis of hyperfine interaction may be

parallel or perpendicular to the apphed field if '] H 1 1 [001], the spectrum is not immediately ob- j vious. However, for //||[111], all hyperfine inter- i actions are equivalent and one sees a nicely re- solved sextet as in figure 5. The probabihty of having \ two ^^Mg ions as neighbors to an anion vacancy ] is about 0.10, and for H\ |[111] one expects a spec- i trum of eleven fines. Some of these are seen in

'

figure 5, and the remainder (except for the central ; hne) can be seen at higher gain. The intensities

| of the hyperfine fines are in accord with expecta- | tion. Hence, the identification of the defect as an i electron in an anion vacancy would seem to be

estabfished. The small hyperfine splittings that are ; observed indicate that the electron is far more ! locafized in the MgO F-center than in those of the ^ alkah hafides. '

Electron-nuclear double resonance (ENDOR)

experiments on F-centers in MgO agree in detail with the expectations for interaction of a nn -^Mg ! nucleus with an electron in an anion vacancy. The

additional information which only ENDOR provides ; is evidence for outward motion of nn Mg nuclei

by 0.04 to 0.08 of the normal Mg-0 distance [11]. i This conclusion results from an interpretation of \ the magnitude of the quadrupole couphng constant; ] ENDOR measurements of the F-center in SrO give ^ similar conclusions [42]. I In the alkali hafides one observes an optical

| band for a defect (F' -center) in which two electrons I are trapped at an anion vacancy. The existence of i a stable F' center in additively colored CaO is

| suggested by an optical band at 3.1 eV and by | bleaching experiments [43, 44]. Irradiation with f fight of energy > 2.8 eV causes bleaching of ! the 3.1 eV band and an enhancement of the F

'

Table 3. ESR parameters of trapped-electron centers ^

F F2 F' excited state F,

g g g D

(cm-')

MgO 2.0023 2.0004 2.0012 2.000 2.0033 0.0287 [47] CaO 2.0001 1.995 1.9980 [35] 1.995 [35] 2.00 [46] SrO 1.9846 1.9816 fev) BaO 1.936

^ Values for which no othfr references are cited are taken from reference [39].

20
x2 1

FiGUBE 5. ESR spectrum ofthe Y-center in MgO.

W||[llll. The central line arises from centers with no "Mg neighbors. The uniformly spaced sextet arises from centers having one nn ^Mg nucleus. The remaining lines belong to a set of eleven Unas from centers having two nn "Mg nuclei.

band at 3.6 eV. There is simultaneously an enhancement of the number of F-centers seen by electron

spin resonance. Irradiation of such samples with

light of energy < 2.5 eV reduces the intensity of

the 3.6 eV F band and enhances that of the 3.1 eV

F' band. It is presumed that the low-energy hght causes release of electrons from shallow traps and

that these are then captured by F-centers to form

F' -centers. The F' -centers do not show ESR absorption in their singlet ground states. However, upon illumination with light of energy > 3.5 eV, a doublet ESR

signal is observed both for MgO and for CaO,

corresponding to an excitation of the triplet state

[35].An F-hke center which apparently has no counterpart in the alkah hahdes is an electron trapped at a

cation-anion vacancy pair (F2-center) [38]. At ~ 9 GHz, its ESR spectrum in MgO, CaO, SrO and

BaO consists of a narrow, but asymmetric, line underlying the central component of the F-center

Une. The F2 line is enhanced at the expense of the

F-line by heating at temperatures up to 900 °C,

followed by :t -irradiation at room temperature to populate the anion vacancy of the vacancy pairs.

After such treatment, no F-centers are detected.

Resolution of parallel and perpendicular ESR hne

components is achieved at 35 GHz, both for MgO

and for CaO; the g'-components are given in table 3

[35]. Unlike the simple F-center, the F2 center may

be depopulated by optical irradiation over a broad

range of energies peaking at 3.5 eV in MgO [45]. A remarkable three-vacancy, two-electron Fcenter has been called the F, center, the subscript

t suggesting that ESR observations are made on

the triplet (excited) state. It appears to be a linear array of a cation vacancy between two anion

vacancies, each having one electron. The axis hes along [001] or equivalent directions [46]. Ex- change couphng of the electrons leads to a singlet

ground state. However, the energy separation Jjk

of the triplet level is reported to be only 80 °K for MgO [47] and 50 °K for CaO [46]. Hence at 300 °K

the triplet levels are thermally populated. For CaO

at 1.6 or 4.2 °K, light in the F-band (3.6 eV) was

found effective in producing polarization of excited

spin levels [46]. The anisotropy of the spectrum

unambiguously identifies the axes. The separation

of electrons is calculated from the zero-field sphtting parameter D to be 4.47 A in MgO and 5.5 A in CaO, assuming D arises whoUy from dipole-dipole

interaction. A central cation vacancy is interposed

to attain stability for the aggregated vacancies.

For MgO, it was found necessary to use doses of

> 2 X 10i« nvt at 150 °C in order to see the F,

spectrum.

5.2. Optical Absorption of Trapped-EIectron

Defects

After a moderate dose of fast-neutron radiation,

MgO exhibits a blue color quite distinct from the

violet appearance after jc -irradiation. It was soon estabhshed that this blue color is unrelated to the

simple F-center. The absorption spectrum of

neutron-irradiated MgO shows an intense peak at 2500 A as well as a large number of bands extending

to 1.2 /X. Unhke the alkali hahdes for which the

electron trapped at an anion vacancy (F-center)

gives rise to absorption in the visible region of the

spectrum, the F-center in MgO absorbs at 2500 A. This has been estabhshed by correlated ESR and

optical absorption studies [15, 48]. This shift to a much higher binding energy for an electron in an oxygen vacancy was predicted remarkably weU

[49]. For MgO crystals which have been heated in magnesium vapor or electron irradiated, there

appears to be a second band very nearly super- imposed upon the F band [50]. Tentatively, this may be interpreted as the F' band (two electrons

in one anion vacancy).

Doubtless the most impressive evidence of

correlation of the 4.9 eV optical band in MgO with

the ESR spectrum comes from the combination of

21
Faraday rotation measurement with ESR observation. There is a marked change in the Faraday

rotation amphtude under conditions of microwave

saturation of the F-center resonance [51, 52]. The optical absorption band of the F-center in CaO is centered at 3350 A [35, 44, 51, 52]. Though

it was not detected at first, a sharp zero-phonon

line at 3555 A appears to be associated with this band [53, 35]. Very recently it has been shown that the Faraday rotation in the zero-phonon line also

is affected by microwave saturation of the electron

spin resonance of the F-center. This appears to be

the first example of a zero-phonon Une associated

with a center of simple symmetry. Its sharpness

and its associated phonon bands should permit

many incisive inquiries of the details of geometry,

of the absorption process and of vibrational modes.

The optical absorption band for SrO is found to peak at 3 eV [35]. It has as yet received little attention.

MgO samples which have been neutron-irradiated

at elevated temperatures or which are heated to 300 °C or more after irradiation show a remarkable

number of bands with sharp zero-phonon transitions. It was predicted that these arise from Faggregate centers [15]. A number of uniaxial stress investigations have been carried out to establish

the symmetry properties of the centers responsible

[54, 55, 56, 45, 35]. A number of these bands are found to be oriented along a [110] direction, as is the two-electron M center in the alkali haiides.

Others have symmetry characteristics resembhng

those of the three-electron R center [45, 35]. Detection and identification of an ESR signal for those

defects in which either the ground state or an

attainable excited state is predicted to be para- magnetic would be most desirable for identifica- tion of the defects. Interaction of an oxide surface with oxygen may

lead to modification of existence defects or generation of new ones upon adsorption. The former case

is typified by MgO, in which the concentration of F-centers in the bulk crystal may be diminished

by neutron irradiation in oxygen. This effect is pronounced if the MgO is in powder form [57, 58]. Destruction of a surface defect (S-center) by O2

is apparently total. Neutron irradiation of alkaline

earth powders in vacuo leads to the S-center de- scribed as an electron trapped at a negative ion

vacancy with five Mg^+ neighbors. An intense color

and the associated ESR spectrum both disappear

rapidly on admission of oxygen. A surface center

labeled S' and having rather different ^-components is generated by UV irradiation of MgO

in vacuum [59]. Admission of oxygen leads to a new defect showing rhombic symmetry. Its gcomponents are ^^^ = 2.077, gyy=-2.00U and

^22 = 2.0073 [60]. These are similar to the values

for O^ observed in KO2 and in alkali haiides

having the O2 defect. The same O^ center is observed on a ZnO surface to which oxygen has been

admitted without UV irradiation.

6. References

[1] W. Low and E. L. Offenbacher in Solid State Physics,

F. Seitz and D. Turnbull, eds., 17, 135 (1965).

[2] J. H. E. Griffiths, J. Owen and I. M. Ward, Repts. Conf. on

Crystalhne Solids, Bristol 1954, p. 81.

[3] P. Auzins and J. E. Wertz, Bull. Am. Phys. Soc. 9, 707 (1964).

[4] J. E. Wertz and P. V. Auzins, J. Phys. Chem. Solids 28,

1557 (1967).

[5] B. Henderson and T. P. P. Hall, Proc. Phys. Soc. 90, 511

(1967).

[6] J. Schneider and O. Schirmer, Z. Naturf. 18A, 20 (1963).

[7] P. H. Kasai, Phys. Rev. 130, 989 (1963).

[8] A. M. Glass and T. M. Searle, J. Chem. Phys. 46, 2092

(1967).

[9] E. R. Lippincott and R. Schroeder, J. Chem. Phys. 23,

1099 (1955).

[10] A. M. Glass, J. Chem. Phys. 46, 2080 (1967).

[11] J. E. Wertz and P. Auzins, Phys. Rev. 106, 484 (1957).

[12] J. H. E. Griffiths and J. W. Orton, Proc. Phys. Soc. (London)

73, 948 (1959).

[13] J. E. Wertz, unpublished data.

[14] J. W. Orton, P. Auzins and J. E. Wertz, Bull. Am. Phys.

Soc. 9 , 739 (1964).

[15] J. E. Wertz, G. S. SaviUe, L. Hall and P. Auzins, Proc.

Brit. Ceram. Soc. 1, 59 (1964).

[16] P. W. KirkUn, P. Auzins and J. E. Wertz, J. Phys. Chem.

SoUds, 26, 1067 (1965).

[17] J. E. Ralph, unpublished reports, AERE. Harwell.

[18] O. Deutschbein, Ann. Phys. 14 , 712 (1932).

[19] A. L. Schawlow, D. L. Wood and A. M. Clogston, Phys.

Rev. Lett. 3 , 544 (1959).

[20] A. L. Schawlow, J. Appl. Phys., Suppl. 33 , 395 (1962).

[21] G. F. Imbusch, A. L. Schawlow, A. D. May and S. Sugano,

Phys. Rev. 140 A 830 (1965).

[22] A. L. Schawlow, A. Piksis and S. Sugano, Phys. Rev. 122,

1469 (1961).

[23] M. D. Sturge, Phys. Rev. 130 , 639 (1963).

[24] J. E. Wertz, P. Auzins, J. H. E. Griffiths and J. W. Orton,

Disc. Faraday Soc. 28, 136 (1959).

[25] F. Seitz, Rev. Mod. Phys. 26, 7 (1954).

[26] Y. Chen and W. A. Sibley, Phys. Rev. 154, 842 (1967).

[27] A. J. Shuskus, J. Chem. Phys. 39, 849 (1963).

[28] J. E. Wertz and P. V. Auzins, Phys. Rev. 139 A 1645 (1965).

[29] F. T. Gamble, R. H. Bartram, C. G. Young, O. R. GiUiam

and P. W. Levy, Phys. Rev. 138 A 577 (1965).

[30] J. H. Hibben, Phys. Rev. 51, 530 (1957).

[31] J. E. Wertz, G. Saville, P. Auzins and J. W. Orton, J. Phys.

Soc. Japan 18, Supp. H, 305 (1963).

[32] R. W. Soshea, A. J. Dekker and J. P. Sturtz, J. Phys. Chem.

Solids, 5, 23 (1958).

[33] J. E. Wertz, J. W. Orton and P. Auzins, J. Appl. Phys.

Suppl. 33, 322 (1962).

[34] D. H. Bowen, Symposium on Mass Transport in Oxides,

Washington, D.C. 1967, p. [35] B. Henderson, Proceedings of the Symposium on Mass

Transport in Oxides, Washington, D.C. (1967).

[36] F. L Gamble, C. G. Bartram, C. G. Young, O. R. GiUiam

and P. W. Levy, Phys. Rev. 134 A 577 (1964).

[37] J. E. Wertz, P. Auzins, R. A. Weeks and R. H. Silsbee, Phys. Rev. 107, 1535 (1957).

22
[38] J. E. Wertz, J. W. Orton and P. Auzins, Disc. Faraday Soc. 30, 40 (1961).

[39] P. Auzins, J. W. Orton and J. E. Wertz, Symposium on Paramagnetic Resonance, I, p. 90, W. Low, ed. (Aca- demic Press, N.Y. 1963).

[40] T. Takeda and A. Watanabe, J. Phys. Soc. Japan 21,

267 (1966).

[41] W. P. Unruh and J. W. Culvahouse, Phys. Rev. 154, 861

(1%7).

[42] J. W. Culvahouse, L. V. Holroyd and J. L. Kolopus, Phys.

Rev. 140 A 1181 (1965).

[43] J. C. Kemp, W. M. Ziniker and E. B. Hensley, to be published.

[44] W. C. Ward and E. B. Hensley, BuU. Am. Phys. Soc. 10,

307 (1965).

[45] R. D. King and B. Henderson, Proc. Brit. Ceram. Soc. 9,

to be published.

[46] D. H. Tanimoto, W. M. Ziniker and J. C. Kemp. Phys. Rev.

Letters, 14, 645 (1%5).

[47] B. Henderson, Brit. J. Appl. Phys. 17 , 851 (1966).

[48] B. Henderson and R. D. King, Phil. Mag. 13, 1149 (1966).

[49] J. C. Kemp and V. L Neeley, Phys. Rev. 132 , 215 (1963).

[50] W. Sibley, to be published.

[51] J. C. Kemp, W. M. Ziniker and J. A. Glaze, Phys. Letters, 22, 37 (1966).

[52] J. C. Kemp, W. M. Ziniker and J. A. Glaze, Proc. Brit. Ceram. Soc. 9, 109 (1967).

[53] J. C. Kemp and W. M. Ziniker, These proceedings.

[54] R. D. JGng and B. Henderson, Proc. Phys. Soc. 89, 153

(1966).

[55] L K. Ludlow and W. A. Runciman, Proc. Phys. Soc. 86,

1081 (1965).

[56] L K. Ludlow, Proc. Phys. Soc. 88 , 763 (1966).

[57] R. L. Nelson and A. J. Tench, J. Chem. Phys. 40, 2736

(1964).

[58] R. L. Nelson, A. J. Tench and B. J. Harmsworth, Trans. Faraday Soc. 63, 1427 (1967).

[59] J. H. Lunsford, J. Phys. Chem. 68 , 2312 (1964).

[60] J. H. Lunsford and J. P. Jayne, J. Chem. Phys. 44, 1487

(1966).

23

Theory of the Energetics of Simple Defects in Oxides

I. M. Boswarva

Queens College, Flushing, New York 11367

and

A. D. Franklin

National Bureau of Standards, Washington, D.C. 20234

Currently, defect formation energy calculations for ionic crystals rest upon the applicability of

(1) two-body short-range interactions, (2) additivity rules that allow assignment of ionic parameters, such as radius, charge, polarizability, etc., independent of ionic displacement, and (3) a simple functional form to express the dependence of the short-range interaction potentials on the distance between

the two ions. An examination of the properties of the alkaline-earth chalcogenides shows that additivity rules for ionic radii and polarizabihties apply as well to them as to the alkah halides. Further, shell- model fits to both the static and dynamic properties of MgO, UO2, and Th02 indicate that a model

based upon two-body potentials can be used to describe these simple oxides. However, attempts to calculate formation energies of vacancies in MgO have been unsuccessful. The problem may lie with

the functional dependence for the oxygen-oxygen interactions, which are attractive in the shell model

fits and are therefore not likely to be well described by the exponential dependence upon ion separation used in the defect calculations.

Key Words: Oxide crystals, point defects, formation energy, theory. Born model, additivity rules.

1. Introduction

The Born model of an ionic solid has been used

rather successfully for the calculation of formation

and motional energies of simple defects in alkah

hedides and the alkahne earth hahdes. Recently,

models such as the shell model have been able to provide a reasonable fit to the lattice dynamic

properties of simple oxides such as MgO [1],' UO2 [2], and Th02 [2]. It therefore seems appropriate to examine the nature of the Born model

calculation of defect energies with respect to its possible apphcation to oxides. In this model, the lattice energy is expressed as

a sum over two-body interaction energies between

all pairs of ions in the crystal. The pair energies

are assumed to depend on the distance between

the ions, the charges of the ions, and their electronic

dipole moments. Parameters describing the properties of the ions, including ionic charge, polarizabihty, van der Waals coefficients, and overlap

energy parameters are needed for the calculation,

and are obtained by fitting the elastic and dielectric properties of the same crystal with the model.

In this paper, we shall first describe briefly the

mechanics of the defect calculations in order to

' Figures in brackets indicate the literature references at the end of this paper.

focus on the basic assumptions involved, then ex- amine the apphcabihty of the concepts of ionic

polarizabihty and radii to a family of compounds

that includes simple oxides (the alkahne earth

chalcogenides), and finally discuss some attempts

at calculating defect formation energies in MgO

and VO2.

2. Defect Formation Energy

Calculations

In the Born model, the crystal is considered to be

static, with separate charged ions situated at positions of equihbrium with respect to the (ionic)

potential energy of the crystal. The electrons are considered to be locaUzed upon particular ions and

to move with the ions. Changes in electron states are included via induced dipole moments, deformation dipoles, etc. In what foUows we will describe

the special case of a vacancy, but the argument

appUes, with trivial modification, to other point

defects. A vacancy is formed by removing an ion from an

otherwise perfect crystal. The zero of energy is thus the energy of the perfect crystal. When the

vacancy is formed, the surrounding ions are dis- placed from their lattice sites and become polarized.

The defect formation energy then contains three

290-621 0-68—

3

25
types of terms: Coulombic, polarization and short- range closed-shell interactions.

Expressions for the first two groups of terms can

readily be written down:

Defect

1 Crystal

I*

J

I, J

eiej

Perfect Crystal

I,

J

ri-rj

where R,, Rj indicate relaxed positions of the ions,

Fi, Fj perfect lattice positions and the summations

are over all ions of the crystals, including defects

in the first term;

m,

where m, is the electronic dipole induced on ion

i and F, [1] is that part of the total field F; at the

ion i due to monopole sources in the crystal, and

is given by

' j^l iRi-Rip \ J M

Further, nii = o;iF,, where a, is the polarizability

of the ion i, and

|Rj-R/|-^

The third group present the major problem, insofar as their form must be obtained from detailed consideration of perfect crystal properties. However,

this difficulty may be separated from the mechanics of the defect calculation and will be discussed

later. For the present, we will assume that the

short range interactions are of the central, two

body type and can be represented by an analytic

form, (f)ij{R). Then

EsR = 'Z {</>o(|R--Rj|)-</)u(|r,-F,|)}.

Because of the short-range nature of these inter- actions only those between close ions, normally

nearest neighbors and next nearest neighbors, are included.

Now the total formation energy

E^Ec + Ei, + E.SR

must be minimized with respect to the ionic relaxations R, — F; to solve the problem. Apart from the

short-range interactions mentioned above and the

problem of obtaining values of polarizabilities a, from perfect lattice data, the only difficulties re- maining are computational ones. The most satisfying way of doing the problem

would be to construct a model crystal containing

the defect plus such a large collection of ions that

all of the important displacements and moments

are explicitly included in the expression for the

total energy. Unfortunately, for ionic crystals this has been possible only for neutral defects. For a charged defect Ec and Ep contain long-range terms,

making the computational problem in such an entirely explicit calculation prohibitively complex.

Instead, the model used consists of an inner region

(I), in which the displacements and moments are treated explicitly, embedded in an outer region (II), treated in a quasi-continuum manner, with the properties of the crystal. The latter is treated in the

harmonic approximation, the total energy of the

crystal depending only quadratically on the ionic

displacements and moments.

Until very recently only calculations in which

region I was just the nearest neighbors of the defect

had been carried out. However, these have been

fairly successful for a number of simple defects, as shown by table 1.

Table 1. Recent defect energy calculations

Formation.

Formation.

Formation. Motion

Formation. Formation. Motion

Formation.

Schottky pair.

Schottky pair. Frenkel pairs and Schottky trios. Vacancy pairs.

Interstitials. Substitutional impurity.

Vacancy.

Vacancies.

NaCl-type alkali halide.

CsCl-type alkali halide. CaF2-type alkaline earth halides. NaCl-type alkali halide. NaCl-type alkali halide. NaCl-type alkali halide. NaCl-type alkah halide. NaCl-type alkali halide.

Boswarva and Lidiard [3]. Tosi and Doyama [4]. Boswarva [5]. FrankHn* [6]. Tharmalingam and Lidiard [7]. Tharmalingam [8]. Tosi and Doyama [4]. Tosi and Doyama [4]. Scholz [9].

*Region I contains nearest and next-nearest neighbors.

26
In general the results are slightly less than the

experimental energies where these are known. But

in this age of large, high-speed computers these

small region I's are unduly restrictive. One of the

authors has just completed a calculation of the

Schottky pair in NaCl extending region I to include

next-neighbors and finds this movement of the I-II boundary typically increases Wg by about 0.3 eV,

from 1.8 eV for a region I containing only nearest neighbors to 2.1 eV for the more extended region I. This large change clearly indicates an unsatisfactory situation. In a detailed examination of the

internal consistency of the model as applied to CaF2, we have found that the electric fields calculated on the basis of the final configuration deviate

significantly from those assumed in the quasicontinuum treatment of region II as far out as the

sixth shell of neighbors to the defect. Calculations

using a more extended region I, such as those now

being done by Scholz at Harwell and Hatcher et al., at Brookhaven, are clearly needed.

The major issues to be resolved in carrying over

this model to the oxides, then, involve the application of two-body short-range interaction potentials

and ionic polarizabilities, and the determination

of suitable parameters describing these from ex- perimental data. It should be noted that once these

issues have been resolved, computer programs de- veloped for the alkali halides can be applied di- rectly to the alkaUne earth oxides, and those for CaF2 to the oxides of the transuranic elements U

and Th, of some of the rare earths such as Ce, and

of the alkali metals.

3. Evaluation of Short-Range

Interionic Potentials

3.1. Semiempirical Approach

It can readily be shown that if the energy £" of a

pair of hydrogen atoms with parallel spins is cal- culated quantum mechanically as a function of

their internuclear distance, r, using first-order perturbation theory and the antisymmetrized product of a pair of Is wave functions, the interaction

is repulsive and the E dependence on r is almost

exponential. A similar dependence has also been

found for the interaction of a pair of He or Ne

atoms. As a result it is often assumed that the

short-range interaction between any pair of closed

shell ions has the form

(/) = B exp {—rip)

where fl, p are constants. Then, if only nearest

'See M. P. Tosi in Solid State Physics, 16, p. 1. Ed., F. P. Seitz and D. Turnbull (Academic Press, N.Y., 19641.

neighbors are included, the lattice energy/ion pair is

W,(r) =Z.^+MB exp (-rip)

where am is the Madelung constant and M is the

number of nearest neighbors that each ion possesses. The parameters B and p can be simply

evaluated, neglecting temperature effects, by the

relations

where V is the volume/ion pair, ro the equilibrium

interionic distance and /3 the compressibihty.

To include next-nearest neighbor short-range

interactions one needs either: (a) to assume some

relations among the six parameters B+-, B+ + , B—, P+-, P++, p—; (b) to employ more experimental information; or (c) to find a reasonable

combination of both. The traditional Born-Mayer

approach, so successful in alkali haUdes, is of the

first kind. It is assumed that Bjj is given by bfiij exp [(r, + r,)/p], where the ionic radii r, and rj are chosen to give as close a fit as possible between

(r+ + r_) and the interionic separations throughout

the family. The Pauling coefficients )8y are given by

/3y = 1 + Zil m + Zj/nj

where Z,, Zj are the valences, and n,, nj the number

of electrons in the outer shell, of the two ions. Under these assumptions, the short-range part of

the perfect-lattice energy can be written as

WsR = Mp+-b exp [(r+ + r_)/p]

+ h M'b[(3+^ exp {2rjp)

+ /3— exp(2r_/p)] exp (-r'/p)

where M' and r' are the number of and distance to the next-nearest neighbors to a given ion, so that

it is still only necessary to evaluate the two parameters, b and p. It can be appreciated that, while

this approach has been fruitful for the alkali halides,

the assumptions involved may not carry over to other substances.

In particular, even though the basic framework

may be applicable to other classes of substances,

there is no guarantee that the same parameters,

including ionic radii and polarizability, need

describe one ion in two different famihes of sub- stances with different structures. Thus when moving to a new substance, for which the same de27
tailed family relationships either don't exist or have not been worked out, one has little alternative

to increasing the amount of experimental data and

decreasing the number of assumptions. This is the

approach taken by one of us in work on CaF2. Still assuming constant hardness parameter p, the four

parameters p, A-\—, A-\-+, A were fitted using

the quantities

where C12, C44, Cn are the elastic constants,

M the reduced mass of an ion pair, €» and the

static and high frequency dielectric constants,

and ojo the circular frequency of the transverse

optic mode. This dispenses with the use of the

additivity rule for ionic radii and the Pauling

relations. Apart from this CaF2 work all the workers quoted

in the table above have used the Born-Mayer

derivation with next nearest neighbors or some

elaboration of it. Frequently, however, the simple

exponential dependence has given too weak an

interaction dependence with respect to the change

in interionic distance and it has been necessary

to include other terms such as ones to represent

the Van der Waals interaction. In this case = Bij exp (- ry/p ) - Cijrjj^ - dij r-^

where the Van der Waals coefficients cy, dij are derived from refractive index data.

3.2. Quantum-Mechanical Approach

In 1948 Lowdin [11] pubhshed the results of a quantum mechanical treatment of the cohesive

energy, interionic distance and elastic constants

of a number of alkali halides. He considered a sys- tem of closed shell ions containing the total number

of electrons A^, for which the Hamiltonian operator,

omitting relativistic effects, is

where Zg is the atomic number of nucleus g, rgg' the

distance between the nuclei g and g' , rig the dis- tance between electron i and nucleus g, r,A- the

distance between electrons i and k, and pi the

momentum of electron i. Neglecting the mutual

deformation of the ions, he used Hartree-Fock

free-ion wave functions. Only short-range inter- actions between nearest-neighbors ions were in- cluded. Even within this framework his task'

presented sizable numerical problems which ne- cessitated several mathematical approximations.

However he found good agreement with experimental values of cohesive energy and lattice parameters for salts with cations and anions of

about the same size. He also found that the three- body interaction arising from part of the last term

of Hop qualitatively accounted for the observed

difference in elastic constants C12 and C44, which

for a model with only two-body effects should be

equal. Even more encouraging to workers using

the Born-Mayer model, he found that the short- range interaction varied exponentially with inter- ionic distance.

Unfortunately, current work by a group of

Lowdin's colleagues at Uppsala (about to be published in Arkiv Fysik 34, 199 (1967)), in which the

required two and three electron overlap integrals are found exactly and next-nearest neighbor inter- actions included, no longer supports the purely

exponential dependence of the short-range inter- actions. Such results strongly reopen the question

of how much the success of the Born-Mayer model

has been fortuitous and certainly indicate that the

first step in the study of a new material such as an

oxide should be at least a limited exploration of

the quantum mechanical aspects.

4. Applicability of the Born Model

to Oxides

The use of any ionic model in defect calculations

rests on the assumption that ionic properties such

as polarizability or short range interaction parameters, which in part are expressible as ionic radii, can be assigned to individual ions and remain un- changed as the ion is displaced from its lattice

site. Very similar to this criterion is the "additivity"

rule, in which unvarying ionic properties are as- signed to ions within a family of substances. It is traditional to use the additivity rule to derive

the properties of alkali and halogen ions from

alkali halide data, and then on the assumption that these ions retain these values, to use the rule to derive properties of other ions from data on their

compounds with the alkali or halogen ions. When this is done to produce parameters for the oxygen ion, it is found that they are not constant

but vary markedly from compound to compound.

This lack of constancy is usually interpreted to mean failure of the oxygen ion generally to follow

"additivity" rules. It seems worthwhile, however,

to explore an alternative possibility. The alkah and 1 halogen ions, while retaining essentially constant

ionic properties within the single family of the rock

salt-type alkali halides, may have different prop28
erties from one family of compounds to another.

A fairer test of the apphcabiHty of the additivity rules to oxides would be to do what was done with

the alkali hahdes — to test them while remaining

within a given family of compounds, all with the same structure. Fortunately, enough of the basic

data now is available for the alkaline earth chal- cogenides to perform such a test.

4.1. Ionic Polarizabilities

Two semi-empirical approaches to the evaluation

of electronic polarizabilities in ionic crystals are

well established. The first, pioneered by Tessman,

Kahn, and Shockley (TKS) [12], assumes (i) that within a family of salts the electronic polarizabilities of each substance is the sum of the electronic

polarizabiUties of the individual ions and that each

ion polarizability is constant within the family.

Under the further assumption (ii) that the electrostatic interactions between ions are purely of the

dipole-dipole type, the effective field acting on an

ion is Eef{ = E + LP where E is the externally applied

field, P the electronic polarization and the Lorentz

factor L is given a value of 477/3 for cubic crystals

with at least tetrahedral site symmetry.

The polarizability per ion pair am is then related

to the refractive index, n, of the salt through the

Clausius-Mossotti relation _ SVm n^-l _

where Vm is the volume occupied by an ion pair.

The polarizabilities ai, are chosen to attain a best

fit to am values throughout the family and the

achievement of such a fit is taken as a measure of

the vaUdity of the assumptions. This fit is found by

minimizing the standard relative deviation

^ m=l ^

with respect to the variables a,, aj when applied

to the family of alkaline earth chalcogenides (13

members included). This analysis produces the

results in table 2.

Table 2. Electronic polarizabilities in (angstroms)^

Mg2+ 0.094 (Fixed) 02- 1.657 1.157 4.497

1.795 Se2- 5.686 Ba2+ 3.188 7.317

It should be noted that the minimum value of F

defines the ionic polarizabilities only to within an

additive constant. Therefore one polarizability must

be fixed by other considerations. In this work the

Mg2+ polarizabihty is taken as 0.094 A'^, the value

found for the gaseous ion by Pauling [13] . However,

this particular ion has a value so much smaller than

the other seven ions in the family that even a 100

percent error in its value would have little effect on the other ions. The minimum value of F is lower

than that obtained with the alkali hahde polarizabilities, indicating excellent adherence to the

additivity rule by the alkahne earth chalcogenides.

These figures have been obtained using the nor- mal Lorentz factor of L = 477/3, which is found by

considering the ions as independent dipoles. If, on the other hand, the charge clouds of neighboring

ions overlapped substantially a smaller value of L

would be appropriate. A plot of different L values

against the minimum of F for each L itself shows a sharp minimum at 477/3 thus giving further evidence

for the appropriateness of the point ion model.

The second approach to finding ionic polarizabilities starts from the relation between the ion

polarizability and the optical absorption due to that ion

d{AE)

where /a(A£') is the absorption coefficient per unit

energy increment A£'; AE = E„' — En, the energy

difference between the two states involved in the

absorption process; A'^ is the number of ions per

unit volume; and C is a material-independent constant. As the absorption due to the cation and the

anion in an ionic sohd are in widely separated

regions of the spectrum this approach should give

polarizabilities directly without the assumptions

needed in the first approach. Unfortunately, the

absorption spectra required he in the far ultraviolet and are not available for most materials (including

the oxides), although Hajj [14] has successfully used

the data of Eby, Teegarden and Dutton [15] to find halide polarizabilities in the alkali halides.

4.2. Van der Waals Coefficients

The equation above for a can be used to obtain van der Waals coefficients. It is assumed that the

bulk of the absorption due to a particular ion occurs near some characteristic energy E so that the

equation reduces to

'Standard relative deviation, F = 2.53 percent. 4>TT^mE

29
where n is the number of electrons in the ion, e the

electronic change, m the electron mass, and h is Planck's constants. With characteristic energies

gotten from the polarizabiUties previously deter- mined the following relations (Mayer [16]) can be

employed to find the dipole-dipole and dipolequadrapole coefficients of the van der Waals inter- action between ions i and j

SEiEjUiaj

2iEi + Ej)

dij 27aiajEiEj

8eHEi + Ej) m nj

Table 3 shows the results obtained.

The only other published values for these oxides,

which are about a factor of 4 smaller, are those

published by Huggins and Sakamoto [17] who

used a method of extrapolating from the constants

for isoelectronic ions in alkali haUdes.

Table 3. Van der Waals coefficients

C__(//s) d^-(A'), d-(A«) d^AA^)

MgO

CaO

SrO

BaO

4.86

45.6

74.2 136

53.0 0.72 41.5 113 329

4.01 49.3

76.5 136

70.6 0.23 34.4

82.9 261

4.3. Ionic Radii

Fumi and Tosi [10] (FT) (1964) have shown that within the framework of the Born model of ionic

solids it is possible to determine completely the

radii of ions within a family of salts using only

soUd state experimental data. This is in contrast

to all earlier evaluations of ionic radii (e.g., Goldschmidt [1]) which, although similarly using the

concept of approximate additivity of interionic

distances, fixed one ion radius either arbitrarily or by appeal to free ion properties. Further, for the alkah hahdes FT found radii for the cations of

about 0.2 A larger than the traditional values with

their anion values correspondingly about 0.2 A

smaller. Work already cited by Boswarva and

Lidiard [3], and Tosi and Doyama [4] found that satisfactory values of characteristic energies of

simple point defects in the alkah hahdes could

only be obtained using FT radii. Therefore it is probably essential to derive ionic radii in the

alkaline earth chalcogenides using the FT approach

for the study of the applicability of the Born model

to defect calculations in this family.

The table 4 gives the results obtained along with

Goldschmidt's values for these ions. The relative standard deviation for the ionic separation cal- culated from these radii compared to observed

values in this family is 6.5 percent compared with

4.4 percent for FT's results for the alkah hahdes.

Bearing in mind (i) the greater uncertainty in ex- perimental data used, especially the compressibihties, and (ii) the greater rehance placed on

the evaluation of ftij through the Pauling relation

for doubly charged ions, it appears that the "additivity" rule in the interionic distance is well

obeyed for this family, another indication of dis- tinct ions in the substances. The good agreement

with Goldschmidt values should be noted so that FT's conclusion for the alkah hahdes that the older

values of the cations are too small by about 0.2 A

and the anion values similarly too large is not mirrored in the alkahne earth chalcogenides.

Table 4. Ionic radii in the alkaline earth chalcogenides,

angstroms

Ours Goldschmidt

0.73 0.78 Ca2+ 1.08 1.06 Si^+ 1.21 1.27 Ba2+ 1.43 1.43 02- 1.40 1.32

1.75 1.74 Se^- 1.87 1.91 Te^- 2.02 2.11

5. Defect Formation Energy

Calculations

5.1. Alkaline Earth Oxides

Yamashita and Kurosawa [19] performed an

early calculation of vacancy formation energies in MgO, CaO, and SrO. They used the TKS values for the ionic polarizabilities, and evaluated the short- range interaction potentials by fitting to cohesive

energy, interionic separation, compressibihty, and

dielectric data. It is interesting to note that they

found the oxygen-oxygen short range interactions

to be attractive, rather than repulsive, thus ruling out any use of the traditional Born-Mayer treatment

involving ionic radii. Furthermore, they found it necessary to assume that the polarizability of the

oxygen ions immediately surrounding the cation

vacancy was a half or less the normal value in order

to give Schottky pair formation energies of reasonable magnitude.

30
We have attempted to use the computer pro- grams, with region I containing the nearest neighbors to the defect, worked out for the alkah halides

to calculate vacancy formation energies in the alkaline earth oxides. The purpose of these very pre- liminary calculations is to explore the effect of

various ways of evaluating the short-range inter- actions. Four methods of obtaining these inter- actions were studied, all assuming exponential

dependence upon the interionic distances. The

resulting Schottky pair formation energies are shown in the table 5, along with results for NaCl

using the same techniques.

Table 5. Schottky pair formation energies, eV

Only 1st n.n. With 2d

n.n. via

ionic radii

Cii, Ci2 and To and - €o + 2

NaCl

MgO

CaO

SrO BaO

1.81 2.1-2.9

2.5 4.5 0.97

1.79 FaUs

0.99

3.58

Fails

1.84 3.971.84

Fails

Column 2 contains the results when only nearest neighbor short-range interactions are included, the

only experimental data needed for evaluating the

parameters being the compressibility /8 and the

interionic distance ro. The value for BaO appears

to be wholly unreahstic and that SrO unreasonably

high compared to the values for MgO and CaO.

For the results in column 3, short-range oxygenoxygen interactions were included by using the

ionic radii and the Pauling coefficients in the Born- Mayer formulation. Again only /8 and ro were needed

to evaluate the short-range interactions. Failure

implies a negative formation energy for the Schottky

pair, the same anamolous result that led Yamashita

and Kurosawa to propose a reduced polarizability

for the oxygen ions immediately neighboring the

cation vacancy.

In column 4 are the results obtained when the

ionic radii were abandoned and the short-range

interaction parameters were determined by fitting

the two elastic constants Cu and Cn, and ro. An

exponential form was assumed for the short-range

potentials for the cation-anion and anion-anion

interactions, with the same hardness parameter

p for each, but different values of the intensity

factor Bij. The cation-cation interactions were

neglected. Only for MgO are the necessary elastic constant data available. Here the oxygen-oxygen

short-range interaction was found to be attractive.

Under these circumstances, there is no justifica- tion for assigning the same values of p to the anionanion as to the cation-anion interaction. To remove

this condition, a fourth experimental equation, that

first deduced by Szigeti [20] involving the dielectric constants and the frequency wo of the transverse

optic mode, was invoked. M is the reduced mass

of the ions and co and e are the static and high

frequency dielectric constants. The result for MgO was again failure of the model to give a positive

Schottky formation energy, again in agreement

with the results of Yamashita and Kurosawa.

It is interesting to note that while the results

for MgO are fluctuating wildly with the method of

evaluating the short-range interactions, the results

for NaCl are very little influenced. This may be

related to the apparent attractive nature found for the oxygen-oxygen interaction. Peckham [1] has

also found an attractive interaction in his shell model fit to both the static and dynamic properties

of MgO. Certainly the assumption of an exponential

functional form for these oxygen interactions must

be highly suspect if they are attractive.

5.2. Uranium Dioxide

Dolling, Cowley, and Woods [2] have obtained neutron scattering data for UO2 and have derived

therefrom the phonon spectra. They have fitted these data with several sets of shell model parameters (see the preceding paper by J. Slater for a

discussion of the shell model). J. D. Axe [21] has

pointed out that some confusion exists as to the

units in which these parameters are presented in the paper of Dolling et al., but that if the shell- shell force constants as given in their table II be

divided by 2 and their ionic charges by V2^ then

the agreement with experiment noted in their

paper can be reproduced. With these corrections,

it is found that their model IV appears to fit the

cohesive energy as well, and to provide a reasonable basis from which to attempt some defect

energy calculations.

Here again, the oxygen-oxygen parameters indi- cate an attractive interionic short-range potential. In the case of UO2 an important defect is the oxygen

interstitial. There is no hope of calculating the

formation energy of the oxygen interstitial with

this attractive oxygen-oxygen potential because

in the interstitial the oxygen ion is enclosed in a cubic cage of lattice oxygen ions, and nothing in the model would prevent the interstitial oxygen ion

from collapsing into one of the lattice ions. Even

for an oxygen vacancy, preliminary attempts to calculate formation energies using a program de- veloped for CaF2 have failed because of inward

collapse of next-nearest neighbor oxygen ions.

31
6. Summary

Defect formation energy calculations for ionic

crystals, as currently being done, rest upon the

apphcability of (1) two-body short-range interionic

potentials, (2) additivity rules that allow assignment (independent of displacement) of parameters,

such as ionic charge, polarizabiUty, and those

determining the short-range interactions, to the

individual ions, and (3) a simple functional form,

such as exponential, to express the dependence

of the short-range interionic potential on the dis- tance between two ions. An examination of the

properties of the alkahne earth chalcogenides as

a family shows that additivity rules for ionic radii and polarizabihties apply as weU to them as to the alkah hahdes. Furthermore, shell-model fits to both the static and dynamic properties of MgO,

UO2, and Th02 indicate that a model based upon

two-body potentials can be used to describe these

simple oxides.

Attempts to calculate vacancy formation energies

in the alkahne earth oxides have, however, been

notably unsuccessful. Since the shell-model

parameters indicate attractive oxygen-oxygen

short-range interactions, it appears hkely that the

problem lies with the choice of the functional

dependence of these interactions upon the interionic

separation. While the exponential form has theoretical backing for closed-shell ions and appears to be vahd for the alkali and alkahne-earth halides,

it probably is not vahd for the oxygen-oxygen

interactions in the oxides. Thus the key problem

in the oxides is to obtain and evaluate expressions

for these interactions.

7. References

[1] G. Peckham, Proc. Phys. Soc. 90, 657 (1967).

[2] G. DoUing, R. A. Cowley and A. D. B. Woods, Can. J. Phys.

43, 1397 (1965) J. D. Axe, BuU. Am. Phys. Soc. 11,

244 (1966).

[3] I. M. Boswarva and A. B. Lidiard, Phil. Mag. 16, 805 (1967).

[4] M. P. Tosi and M. Doyama, Phys. Rev. 151, 642 (1966).

[5] I. M. Boswarva, Phil. Mag. 16, 827 (1967).

[6] A. D. Franklin, Submitted to J. Phys. Chem. Sol.

[7] K. Tharmahngam and A. B. Lidiard, Phil. Mag. 6, 1157

(1961).

[8] K. Tharmalingam, J. Phys. Chem. Sol. 24, 1380 (1963);

Ibid, 25, 255 (1964).

[9] A. Scholz, Phys. Stat. Sol. 7, 973 (1964).

[10] F. G. Fumi and M. P. Tosi, J. Phys. Chem. Sol. 25, 31

(1964); M. P. Tosi and F. G. Fumi, Ibid 25 , 45 (1964).

[11] P. Lowdin, A Theoretical Investigation into Some Properties of Ionic Crystals (Uppsala, 1948).

[12] J. R. Tessman, A. H. Kahn and W. Shockley, Phys. Rev. 92, 890 (1953).

[13] L. Pauhng, Proc. Roy. Soc. A114, 181 (1927).

[14] F. Hajj, J. Chem. Phys. 44, 4618 (1966).

[15] J. E. Eby, K. J. Teegarden and D. B. Dutton, Phys. Rev. 116, 1099 (1959).

[16] J. E. Mayer, J. Chem. Phys. 1, 270 (1933).

[17] M. L. Huggins and Y. Sakamoto, J. Phys. Soc. Jap. 12,

241 (1957).

[18] V. M. Goldschmidt, Skrifter det Norske Videnskaps Akad- emie No. 2 (1926), No. 8 (1927).

[19] J. Yamashita and T. Kurosawa, J. Phys. Soc. Jap. 9, 944

(1954).

[20] B. Szigeti, Proc. Roy. Soc. A, 204, 52 (1950).

32
Lattice Dynamics and Ionic Interactions

John Slater, J. J. Thomson

Physical Laboratory,

Whiteknights Park, Reading, England

A ver>' brief review is made of topics relating to the use of models of forces between atoms in insulating cr>'stals. The topics include elastic constants, lattice dynamics, lattice energy, pair potentials and ionic radii. Key Words: Elastic constants, ionic radii, lattice dynamics, lattice energy, pair potentials.

1. Introduction

The title of this review allows a great deal of scope, and no single topic will be dealt with in much detail. The question to be answered is roughly,

"What sorts of models can be used for the forces

between atoms in insulators?" The point of view

will be that of a physicist who knows httle about mass transport: the aim will be to reiterate things

that are "weU-known" to provide a background

for subsequent papers. Work on this subject has

progressed along a few rather distinct Unes, some

offering more detail than others. The headings

wUl be — lattice vibrations and force constants; anharmonicity; elastic constants; defect appUcations; lattice energy; ionic radii, and miscellaneous

topics.

2. Lattice Vibrations and Force

Constants

The dynamics of crystal lattices is usually dis- cussed [1] ' in terms of the crystal potential energy

expanded about the equilibrium positions of the

ions in the ionic displacements Ua ^^j. Small vibrations of the lattice are then controlled by the

second derivatives

0 G-')lo'

If we ignore the higher terms we have the "harmonic" approximation — the forces are linear in the

displacements — the "springs" are Hookean. The

' Figures in brackets indicate the literature references at the end of this paper.

above derivative is the force on the ion (

. ) when

//'\

the ion L , j is displaced a unit amount, all other

ions being fixed. It is usually called an interatomic

force constant and work in lattice dynamics aims

to determine these from experimental observations on lattice waves. The experimental data which have to be explained

are several: thermal effects, for example conductivity, infrared absorption, Raman effect and especially x-ray diffuse scattering and neutron inelastic scattering. The last is the most important. It is convenient to deal with plane waves of propagation

vector k. If there are n atoms per unit cell in the

crystal there will be 3n kinds of wave with a given

k and we want to know the frequencies a; asso- ciated with each of these. It is clear that figure 1, which shows the experimental dispersion curves

for UO2 determined by neutron diffraction [2]

,

contains a great deal of impUcit information about

the crystal forces. The procedure used is to assume some model for the forces with suitable parameters

and then either fit these by least squares to the ex- perimental dispersion curves (as in fig. 1) or deter- mine the parameters from macroscopic constants,

and then see how the agreement with the neutron

results looks. The simplest model used is the "rigid-ion"

model [3], with point-ion long-range Coulomb

forces and arbitrary short-range forces, usually

restricted to first or second neighbours. The comparison with experiment of a nearest neighbour

rigid-ion treatment [4] of Nal is shown in figure 2. For a crystal Hke germanium, this would imply

short-range forces only, and this approach was

found to be incorrect: it emerged that the polarisation of the ions had to be introduced. This gave

rise to the "deformation-dipole" and "shell" models, which are essentially equivalent [5]. In the

33
shell model one supposes that the outer electrons

of an ion can be represented as a rigid shell attached

to the nucleus and core by an isotropic spring.

The polarisation of an ion is then represented by

the relative displacement of shell and core. Figure

2 shows the improvement this gives for Nal.

Both models, however, give a bad fit for the

longitudinal optic branch (highest branch) at the

(111) zone boundary. This point corresponds to a wave in which the anions are essentially stationary

and alternate planes of cations move in opposite

directions [6]. The rigidity of the anion shell seems

to be raising the frequency. The so-called "breathing-shell" model introduced by Schroder and co- workers [7, 8] allows a compressible shell and

appears to give significant improvement over the

simple shell model for some substances (for ex- ample, Nal) without introducing more parameters.

All of these models are classical, although they can be given a quantum-mechanical basis. For MgO, Peckham [9] used a shell model fitted to the elastic and dielectric constants: the agree- ment is shown in figure 3. The fit here was not improved significantly by the introduction of a breathing shell model. The number of oxides which

have been investigated is small: MgO, UO2, Th02,

a few branches in quartz [10] and some work on

crystals like MnO and NiO, where the main interest

is in magnetic properties. It is expected that CaO

will be treated in the near future. There are three reasons for this lack of information: the experiments are slow and expensive, it is difficult to obtain good crystals of sufficient size, and difficult to observe high-frequency optic branches where

phonon populations are low at room temperature.

[OAO] [O.O.l]-[M,0] - C [0.0.0] t

- REDUCED WAVE VECTOR COORDINATE. ^

Figure 2. Phonon dispersion in sodium iodide [4]. Reproduced from Phonons p. 66 Ed. R. W. H. Stevenson, Oliver and Boyd, Edin- burgh (1966).

3. Anharmonic Contributions to the

Crystal Forces

In the perfect crystal lattice, anharmonic terms

hke the third derivatives

34
give rise to macroscopic effects such as thermal

expansion and the temperature dependence of

elastic and dielectric constants, while on the quasiparticle model, anharmonicity accounts for the short

life times of phonons, observed in thermal conductivity and neutron scattering. The number of short

range anharmonic constants allowed by symmetry

is large and, in planning calculations, simphfications have to be made, usually an assumption of

axiaUy symmetric nearest neighbour interactions or

a Born-Mayer potential. The Coulomb part of the

anharmonicity is ignored. Nevertheless these fairly

crude models have been quite successful in explaining, for example, thermal expansion [11] and the

behaviour of neutron groups with change in temperature [12, 13]. The number of calculations

which have been carried out is relatively small.

4. Elastic Constants

The extent of the deviation of a crystal from the

radial potential type of behaviour is traditionally

estimated from the agreement of the elastic constants with the Cauchy relations. For crystals with

complex structures, hke Si02, one can hardly ex- pect these to be satisfied, and indeed they are not. For an inert gas configuration like MgO, one might

hope for better, but the values given in table 1 are

not encouraging. On the other hand it is not clear

that the deviation from the Cauchy relations implies

any form of covalent bonding: for example, it ap- pears that Schroder's "breathing-shell" model,

mentioned in section 2, gives a deviation even if only central forces are allowed.

Table 1. Elastic constants (dyn. cm~^XiO") of some crystals

in the sodium chloride structure

Showing the deviation from the Cauchy relations.

Cl2 C44/C12

NaCl 1.23 1.26 1.03 LiF 4.04 5.54 1.37 MgO 8.7 14.8 1.70

5. Defect Applications

It is clear from figure 3 and others that the agree- ment between theory and experiment for lattice vibrations can be very close. It would be a pity

then, even in the absence of a complete physical

justification for the forces involved, not to use this information for something, and a very convenient

by-product for some defect applications is Green'sfunction information. Figure 4 shows what is essentially the mechanical impedance of a magnesium

ion in an MgO crystal. What is convenient is that one can convert this in a trivial manner to the case where an atom of different mass is substituted, or the force constants around the defect changed.

This sort of function is particularly convenient for comparison with impurity infrared absorption and

the vibrational side-band structure on optical ab- sorption lines of impurity ions. The rich detail of

these together with the reliability of the lattice dynamical calculations permits information to be

obtained about the defect environment. For ex- ample, it can be established from the sideband

structure that V'^+ and Ni^^ substituted in the MgO

.rx^..,

^. \ \ \ \ \

(E.C.o) 1 (C.E.P

Figure 3. Phonon dispersion in magnesia.

A composite of points measured by Peckham [9] and Buckland and Saunderson (29]. The curve is the theoretical shell-model prediction of Peckham.

35
Figure 4. The mechanical admittance for radial (Aig) motion of the anions surrounding a magnesium

ion in a magnesium oxide crystal.

lattice have greater force constants to neighbours

than the Mg^+ ion itself [14].

6. Interatomic Potentials and

Lattice Energy

Various potential functions have been used to describe the repulsive forces which balance the

Coulombic attraction of the ions in an ionic crystal.

These repulsive forces arise essentially from the

Pauli principle — from the increase in electron

kinetic energy as the electron density in the crystal

increases.

a. For rare gas solids the Lennard-Jones potential is usually used.

V(r)=-^ + ^n (1)

the first term being the van der Waals attractive

potential; n is usually about 12. Exponential and

polynomial two-part potentials have also been used,

and it is hard to find solid arguments for preferring one to another [15].

One reason for mentioning work on the inert gases is that it is probably here that the great

progress made recently in formal quantum mechanical theory of atoms will first show results for sohds. Several calculations have been made using

variational methods of the interaction of inert gas atoms; for example D. Y. Kim [16] has calculated

the potential for two interacting He ions, obtaining

the following analytic form (r in A, V in eV):

r(r) =387e-4-42'- + 366e-^>5'--0.58/A^-«. (2)

The second exponential term, which becomes important for r ^ 3 A, is worthy of note, even if helium

is a rather special case.

b. For ionic crystals in the sodium chloride structure the time-honoured form is the Born-Mayer

potential [1]

V{r)=-y+Be-'-lp. (3)

The exponential dependence can be justified by

consideration of the falUng off with distance of

the electronic charge density.

c. The most generafised potential in common use

is obtained by adding van der Waals terms to the

Born-Mayer potential

Vir)=j + Be-'-lp-ClT^-Dlt^. (4)

This potential has been used recently by Ladd and

Lee [17] for the discussion of lattice energies in a wide range of ionic crystals, including oxides.

36
d. There are miscellaneous other potentials used

of logarithmic and polynomial form in 1/r for which

it is hard to find physical justification. One can also regard the Born-Mayer-Verwey [30] potential as rather unphysical: it is essentially a Born-Mayer

potential in which the forces are strengthened inside

the equilibrium distance. Perhaps another exponential term as in eq (2) would be preferable.

It is worth emphasising that these potentials

apply to simple structures in the electrostatic

model. In more complex structures direct polarisation can play a very important part. For example in the spinel structure oxygen ion polarisation is estimated to give 25 percent of the cohesive energy

[18]. Also the use of the electrostatic model implies

that one can associate a particular electron with a

particular atom and treat the forces then as arising

from the interaction of charge distributions on these

distinct ions. If one is deahng with a crystal hke

MgO which may approximate a closed-sheU configuration then one can perhaps hope that this will work, and the success of the shell model for the

lattice dynamics of MgO offers some evidence

that it does.

When one considers oxides which are more inter- esting in practice, the oxides of transition metals,

for example, the electrostatic model can teU only

part of the story. Molecular orbital theory has had

considerably more success in explaining the spectra

of transition metal ions in crystals than the older,

essentially electrostatic crystal field theory, and

one must expect that work on ion motion will have

to rely upon this sort of theory, at least in a qualitative way, if it is to give results. Quite extensive

work of this kind has been done on the wide range

of materials of the spinel structure [18, 19].

7. Ionic Radii

So far the impression has perhaps been given

that the situation in interionic forces is unchanged

since thirty years ago, and to redress the balance,

something should be said about recent moves away

from the older forms. A. R. Ruffa [20] has worked

on the polarisabilities and susceptibilities of ions

Ij in crystals. On the basis of quantum mechanical

I arguments of an approximative kind, he has shown

jl that the additivity rule for the polarisabihties of

the ions in the alkali haUdes is something of an

i accident, and suggests that the cation polarisability

; may be considerably larger than the free ion value. He has used related arguments for rutile polarisability and in explaining the magnetic susceptibilities

of the alkah halides.

It is now some time since Tosi and Fumi [21] pubj lished their ionic radii, but it is perhaps worth reiterating their conclusions, because they seem very

important. The use of ionic radii for alkali halides is justified because of the success of the additivity

in prediction of interatomic spacings in the crystals.

However the additivity cannot alone provide a unique set of ionic radii. One more condition is necessary. For the Pauling and Goldschmidt radii this is obtained in effect from the polarisability of the

ions in solution. It can be argued that this gives a

free ion radius rather than a crystal ion radius,

and Tosi and Fumi have calculated ionic radii from

crystal data alone using the assumption of an inter- atomic potential of the Born-Mayer type. They have

obtained in this manner radii which differ by

amounts between 0.15 and 0.3 A from the tradi- tional values, the cations being larger and the anions

smaller than in the traditional case. Table 2 shows

an example. The radii agree well with those ob- tained from x-ray measurements of electron density.

The radius of a given ion varies somewhat from one

crystal to another.

Table 2. Comparison of ionic radii

Ion Traditional Radius of Fumi

radius and Tosi

Na+ 0.95-0.98 A 1.21 A

ci- 1.81 A 1.65 A

It has always been the philosophy, of course, to use ionic radii with some caution. Wyckoff writes

in the introduction to his tables of crystal structures: "We nowadays talk about ionic, covalent and

other bond 'sizes' with an assurance that may

exceed our knowledge of the physical realities of

the situation." One supposes that the radii tra- ditionally used are to be questioned also for less simple ions. J. C. Slater has recently reviewed the

subject of radii obtained by these older methods [25].

8. Miscellaneous Topics

A great deal of work of a quantum mechanical

kind has been done falling roughly within the scope

of this talk, which the writer has neither the space

nor the authority to deal with. Four of these are mentioned here.

a. Quantum mechanical calculations have been

carried out by several Scandinavian workers on

non-stoichiometric oxides like zirconia and praseo- dymia [22, 23].
b. A large amount of detailed work has been done

by physical chemists on molecules like MgO, SrO

etc. Recently Tosi and Doyama [24] have applied

a model borrowed from lattice dynamics to treat the alkali hahde molecules and obtain good results. One supposes that some flow of information is possible in the other direction.

c. The fact that the free ion does not exist has inspired calculations of form factors and

electronic states of oxygen ions and the associated

cations in crystals Uke MnO and MgO [27, 28].

d. The work of Benneman, Kleinman and Phillips

should also be mentioned as offering the promise

of a better understanding of binding in insulating

crystals [26]. In summary, it seems that for a crystal of MgO

(inert gas) type the Tosi and Fumi radii and BornMayer potential, with Van der Waals terms if these can be estimated, offer as reliable a model as is available, though this simple model may well turn

out to be mistaken (see remarks by Boswarva and

Franklin). For more complex structures, the situa- tion is more uncertain, although all is not yet lost; polarisation energies and bonding energies can be

estimated; Madelung energies can be computed

trivially. Great progress is being made in the detailed relation of theory and experiment in the field of lattice dynamics and thermal properties for simple crystals. With regard to other perfect lattice properties that have been mentioned, work is being

done on many fronts, and it is encouraging to note

that much of the progress is comparatively recent. With the great strides that have been made recently

both in computing techniques and in formal quantum mechanics there is reason to look forward to detailed fundamental treatments of the nature of

forces in crystals.

9. References

[1] M. Born and K. Huang, Dynamical Theory of Crystal

Lattices (Oxford University Press, 1954).

[2] G. Dolling, R. A. Cowley and A. D. B. Woods, Can. J. Phys. 42, 1397 (1965).

[3] E. W. Kellerman, Phil. Trans. Roy. Soc. A 2 38, 513 (1940).

[4] W. Cochran, Phonons Ed. R. W. H. Stevenson, Ohver

and Boyd, Edinburgh (1966).

[5] R. A. Cowley, Proc. Roy. Soc. A 268, 109 (1962).

[6] A. D. B. Woods, B. N. Brockhouse, R. A. Cowley and W.

Cochran, Phys. Rev. 131, 1038 (1963).

[7] U. Schroder, Solid State Commun. 4, 347 (1966).

[8] V. Nusslein and U. Schroder, Phys. Stat. Sol. 21, 309 (1967).

[9] G. Peckham, Proc. Phys. Soc. 90, 657 (1967).

[10] M. M. Elcombe, Proc. Phys. Soc. 91, 947 (1967).

[11] R. A. Cowley, Proc. Phys. Soc. 88, 473 (1966).

[12] R. A. Cowley, Phonons Ed. R. W. H. Stevenson, Oliver and Boyd, Edinburgh (1966).

[13] B. Dawson. A. C. Hurley and V. W. Maslen, Proc. Roy.

Soc. A. 298, 289 (1967).

[14] M. J. L. Sangster, Reading University (unpublished).

[15] G. L. PoUack, Revs. Mod. Phys. 36, 765 (1964).

[16] D. Y. Kim, Revs. Mod. Phys. 35 , 443 (1963).

[17] M. F. C. Ladd and W. H. Lee, Progress in Solid State Chemistry, Ed. Reiss. 1 (1964) and 2 (Pergamon Press,

Inc., New York, N.Y., 1965).

[18] G. Blasse, Phil. Res. Repts. Suppl. (1963-4).

[19] S. lida and H. Miwa, J. Phys. Soc. Japan 21, 2505 (1966).

[20] A. R. Ruflfa, Phys. Rev. 159, 742 (1967).

[21] F. G. Fumi and M. P. Tosi, J. Phys. Chem. Solids 25,

31 (1964).

[22] B. Roos, Ark. Fys. 25 , 363 (1963).

[23] B. G. Hyde, D. J. M. Bevan and L. Eyring, Phil. Trans. Roy.

Soc. 259, 583 (1%6).

[24] M. P. Tosi and M. Doyama, Phys. Rev. 160 , 716 (1967).

[25] J. C. Slater, J. Chem. Phys. 41, 3199 (1964).

[26] J. C. PhiUips, Phys. Rev. Letters 19, 415 (1967).

[27] S. Nagai, J. Phys. Soc. Japan 22, 457 (1967).

[28] P. M. Raccah and R. J. Arnott, Phys. Rev. 153, 1028 (1967).

[29] R. J. Buckland and D. H. Saunderson, Harwell Research

Report A.E.R.E.-R5467 (1967).

[30] R. Guccione, M. P. Tosi and M. Asdente, J. Phys. Chem.

Solids 10, 162 (1959).

38
On Electron-Lattice Interaction in the CaO F Center ^

James C. Kemp and William M. Ziniker^

Department of Physics, University of Oregon, Eugene, Oregon 97403

The sharp vibronic structure in several absorption

bands of point defects in alkaline earth oxides has

gotten considerable study, because of the relation- ship to the phonon spectrum. Unhappily, none of

the vibronic structure analyzed to date has been

identified with any very simple point defect, such

as the F center, nor with any center of estabhshed

structure. Thus electron-phonon coupling coeffi- cients to specific lattice modes could not be seri- ously estimated; nor could local modes be discussed.

Recently we identified the zero-phonon peak ^ (3555 A) and associated band ^ (center 3350 A) of

the CaO F center, by the method of microwavesensitive Faraday rotation. One has an opportunity

with the simple electronic structure involved here

to understand the optical, vibronic, and magnetooptic properties in detail. The spectral Faraday

rotation, d{E), for example, seems to have an

obvious interpretation. The observed overall band

rotation pattern is quite analogous to that in alkali haUde F centers. On the other hand the d(E) struc- ture accompanying the zero-phonon peak is absorption derivative-like, i.e., dispersion-like, with (in particular) a positive lower-energy peak. The zero- phonon transition is evidently to an orbital singlet,

indicating a large Jahn-TeUer splitting of the ex- cited p state, so that orbital magnetism is quenched.

We invoke non-cubic modes, such as a local Tag

mode of the first cation shell. The associated 6{E)

should then be simply — A./2e times the spectral

dispersion angle of the zero-phonon peak, where

A. is the spin-orbit constant and e is a (mean) Jahn-

' Research supported by the U.S. Air Force Office of Scientific Research.

^ Present address: Ceramic Engineering Department, University of Washington,

Seattle, Washington.

= W. M. Ziniker, BuU. Am. Phys. Soc, Ser. II. 12, 896 (19671. 'J. C. Kemp, W. M. Ziniker, and J. A. Glaze, Proc. Brit. Cer. Soc. 9, 109 (1967).

Teller sphtting. (This pertains to the paramagnetic

6 seen at hehum temperatures.) From moments of

the overall F-band magneto-optic pattern we know

X = — (8± 1) X 10~^ eV. The observed zero-phonon

d{E) curve is correctly fitted, where we use

e = 0.15 eV — a reasonable value if strong noncubic

relaxation is assumed. A picture using a noncubic

configuration coordinate (added to interaction with

other modes such as the familiar breathing mode)

is thus consistent with the rotation patterns of

both the broad band and the zero-phonon line. The

negative \ needed incidentally strongly supports our identification of the F-center optical spectrum

in CaO: Another species which we have suspected

might have many properties similar to those of the

F center is interstitial Ca* +; but this would doubtless

have positive k in the p state. Other structure on the CaO F band consists of

peaks at 3531 A and 3518 A and some weaker

peaks, including some possible barely resolved

weak structure just to the high energy side of the

zero phonon line. The interval 3555 to 3531 A is about four times the CaO reststrahl energy of

0.045 eV. Actual comparison with the lattice spec- trum cannot yet be reported.

We wish to speculate that a model in terms of

bulk lattice modes may not be appropriate in alkaline earth oxide F centers. The degree of local distortion would seem to favor strong local modes,

electron lattice coupling being then indirect, via a broadening of the local modes. The latter can be

treated theoretically in this structure using point

ion-lattice techniques and simple F-electron wave

functions, and the frequencies and couplings pre- dicted for at least two or three of the simplest

modes. (The Tig mode has already been dealt with

in MgO, etc., in different contexts.)

39

Electronic States of Defects in Irradiated Oxides

B. Henderson

Physics Dept., Keele University, Keele, England

This paper reviews some of the progress made during the last few years in understanding the

electronic states of defects in irradiated oxides. The discussion refers particularly to the alkaline earth oxides. A short introduction indicating the scope of the work on oxide materials, is followed by

a description of the main types of experiment proving useful in the simple cubic oxides. In particular the following topics are discussed: (i) the electron-phonon interaction in the MgO F-centre, (ii) the evidence for the F' -centres in the oxides, (iii) the anneahng of radiation damage and the production

of aggregate centres, (iv) the investigation of the structure of defects in oxides using uniaxial stress

splitting of optical zero-phonon lines. Some novel results of the e.p.r. spectrum of a triplet (S=l) state defect in neutron irradiated AI2O3 are also discussed: a new model is proposed for this defect.

Key Words: Annesding, defects, electron-phonon, electronic states, interaction, radiation damage.

1. Introduction

Irradiation of insulating solids may produce

varied changes in their physical properties conse- quent upon the introduction of defects of atomic

dimensions into the lattices. These defects, which

trap electrons or holes, may be investigated by the

techniques of optical absorption and/or electron

spin resonance. Such techniques have been

especially fruitful in the alkali halides, and in these compounds the electronic properties of many

defects (F-centres, F-aggregate centres and V- centres) are well understood. Similar studies on

oxide systems form a natural extension of this work.

However, the additional trapping potential of the

oxygen ion vacancy as well as the variety of crystal structures in which it might be found provide

complexities which are not normally met with in the alkali haUdes. In addition, most oxides are

relatively resistant to radiation damage by ionising

radiation so that reactor irradiation is frequently

necessary, with the result that defect sites are pro- duced on both the anion and cation sublattices.

Unfortunately, most commercially available oxides

are relatively impure compared with the alkali halides, and consequently great care has to be taken

to recognise effects due to impurity ions. UnUke the alkali halides, oxide systems have been

the object of few serious optical studies and the

absorption spectrum of many defects is completely

unknown. Thus for many defects the only source

' Figures in brackets indicate the literature references at the end of this paper.

of identification is their paramagnetic resonance

spectrum. Many such spectra in oxides have been

attributed to electrons trapped at oxygen ion va- cancies. The paramagnetic states of these defects

will have their orbital momentum quenched and

thus their g'-values are close to the free electron

value. In the alkaUne earth oxides there is a small

negative shift from the free electron ^-value [1];^ a somewhat larger shift is observed in barium

titanate [2]. In alumina [3], calcium tungstate and

lithium niobate [4], and in numerous vitrous ox- ides [5], significant positive gr-shifts are observed.

This subject will not be discussed further except

to note that no satisfactory theory exists for the

g-value of electrons trapped at oxygen ion vacancies.

Instead the discussion will be centred on the recognition of oxygen ion vacancies, and vacancy aggre- gates in the alkaline earth oxides, and of a trapped

hole centre in AI2O3.

2. Single Vacancies in the

Alkaline Earth Oxides

2.1. The F-Band in MgO and CaO

The spectral position of the F-band in MgO

(4.9 eV), CaO (3.65 eV) and SrO (3.0 eV) has now

been determined with some certainty. The elec- tron bound to the anion vacancy is in many ways

analogous to a hydrogen atom, and the localised

electronic states of the F-centre are conveniently

described by a /s (Ai) ground state and a 2p (Fi)

excited state. An electric dipole transition between

these levels gives rise to a broad bell-shaped

290-621 0-68—4 41
absorption band, whose width and shape reflects the strength of the coupling between the electronic

states of the centre and vibrational modes of the

lattice. In the configurational coordinate model,

which we use here, it is assumed that the coupHng

is to only one normal mode of the lattice. The model

predicts that the half-width (//o) at 0 °K is related

to the vibrational frequency Vg by

where S is a vibrationless quantity representing the

effective number of phonons excited in the transition. The breadth of the band H{T} is then related

to the absolute temperature {T) by

H{TY = H{OY co\h^- (2)

The peak position of the band also shifts with temperature. This shift is produced by the combined

effects of the change in lattice constant with temperature and the change in force constants of the

oscillators in the ground and excited states. The

shift consequent upon the variation in lattice parameter may be estimated from the thermal expansion

data and by assuming an Ivey-MoUwo Law to apply

to the oxide F-bands as in the alkah hahdes. The

change produced by the difference in force constants may be estimated by applying the method

of moments to the band shape. Applying this to the

first moment we see that the shift of the band

centroid to lower frequencies as a function of temperature is

•

: LE^hfhvg coi\y^ (3)

where / is the fractional change in force constants [6]. In terms of the Vg and Ve, which refer

to the vibrational frequencies interacting with the

lattice in the ground and excited states, respectively.

If we have sufficient experimental data to deter- mine S, Vg and /, we can construct the configurational coordinate diagram for the ground and excited

states of the F-centre by methods of Klick et al. [7]. The experimental data, H{T) and E{T) for the F-centre in MgO are shown in figure 1. The half- width can be fitted accurately to a curve of the

form (2) with 1/9=7.8x10125-1. Thus the HuangRhys factor determined by substitution for Vg in

(1) is S — Z9. Consideration of the peak position

0 5 10 15 20

(T°K)'''2

Figure 1. The experimentally determined F-band peak E(T)

and half-width H(T) plotted as a function o/T"'^.

The solid line through the half-width data is the curve H(T)^ = {H(0)^ coth where i'j=7.8x 10'^ sec"'. This frequency occurs just below a peak in the experi- mental phonon density of states curve.

temperature dependence leads to Ve~7AX W^S~^.

The most important predictions based upon the

configurational coordinate curves which may be

constructed from this data relate to the probability of observing luminescent emission from the F-centre excited state. Consideration of the cri- terion for luminescence discussed by Dexter et

al. [8] , suggests that in the present case, the re- turn of the F-centre electron to the ground state should be via a luminescent transition of high probability at a photon energy of about 3.0 eV. Such a

transition has not so far been observed in MgO

F-centres.

It is of some interest to note the relatively large

value of the Huang-Rhys factor for the F-centre

in MgO (S = 39). In CaO, Bennett has calculated

that S = 3.6 which agrees well with the value pre- dicted from the zero-temperature half-width as- suming that the vibrational frequencies coupled

to the F-centre in the ground state are the fundamental longitudinal optic modes. Now at r=0 °K

the probabihty of observing a transition involving

n-vibrational quanta is then given by [9]

42
Thus when Pn is normaUzed to unity, the transition

probabiUty of the zero-phonon Hne relative to the

whole band is e'^. Thus, for the F-centre in MgO,

the zero-phonon line transition probability is very

small (since S = 39). However, for CaO the zero- phonon hne has a quite reasonable probabihty since

S is of order only — 6. These ideas seem to be in agreement with experimental results since no struc- ture is associated with the F-band in MgO whereas

in CaO zero-phonon lines have been observed in the appropriate position with respect to the F-band

in both absorption and emission. The recent measurements of the zero-phonon line under uniaxial

stress apphed along [100], [110] and [111] axes show that the transition is A -^T within a centre with cubic symmetry [10]. The stress sphtting

pattern for such a centre is simple and unmistakable [8, 9]. Thus this assignment is contrary to the report by Kemp and Ziniker (this conference)

that this zero-phonon transition, in the presence of

strong Jahn-Teller interaction is to an orbital

singlet. Further experimental evidence is required

to reconcile these data and to confirm^ (or other- wise) that the zero-phonon hne at 3557.3 A is indeed

due to the F-centre.

2.2. F'-centres in MgO and CaO

In the alkahne earth oxides, the F' -centre is an oxygen ion vacancy containing two trapped electrons. Unlike the F-centre it is, therefore, electrically

neutral and would be the stable anion defect in an

otherwise perfect crystal. Studies of these centres

are few. Ward and Hensley [11] have identified an

optical band at 3.1 eV with the F'-centre in CaO

and Kemp et al. [12] have observed interconversion

between the F and F'-centres in CaO using photoexcitation techniques. We have recently observed

spin resonance spectra under ultra-violet illumination which we believe to be due to excited spin

triplet states (S=l) within the helium-hke F'- centres in MgO and CaO. The MgO specimens

used in the experiments had been annealed in •hydrogen at 1600 °C prior to irradiation to doses

in the range 2-20xl0i8 nvt. The CaO single

crystals used were taken from the region of the

molten pool closest to the graphite electrodes; the

reducing effect of the graphite produced excess Ca in the samples as evidenced by the copious

numbers of F and F'-centres apparent from the

optical absorption spectrum.

The bound electronic states of the F'-centre are expected to be similar to the states of the two

electrons bound to the radial potential of a hehum

hke atom. The ground state is a spin singlet 'S state

in which both electrons are in the lowest energy

2p y

Is 25

3 lev

allowed

forbidden

Figure 2. Possible energy of levels of the helium-like F'-centres:

in absorption only 'S —* 'P is allowed.

spatial state (Is). The excited states correspond

to one of the electrons having spatial symmetry

appropriate to 2s or 2p states. Figure 2 shows that the next states in order of increasing energy will be ^S, 'S, ^P, and 'P. The only allowed optical

transition is then from the ground state to the 'P state. This transition occurs at 3.1 eV in CaO

[11] although the process by which the triplet state becomes occupied in both CaO and MgO is not

yet understood. In figure 3 the photo-excited reso- nance observed in CaO is shown. A similar reso- nance is observed in MgO with a clearly defined

hyperfine structure associated with Mg^^ attending

the central single hne. The g-values are respectively

1.9866 for CaO and 2.000 for MgO. It should be

noted that in the case of MgO all hght below 3.5 eV

was filtered out thus giving a lower bound for the

^P transition energy. The CaO photo-excitation was accompanied by a strong, long-hved red

luminescence at room temperature. The evidence,

although not conclusive, strongly supports the

interpretation that these photo-excited resonances

are due to excited ^S-states of the F'-centre. Further

studies of these resonances are in progress.

3. Annealing Studies

3.1. F2-Centres

When neutron irradiated MgO is annealed at temperatures above 300 °C, changes in the e.p.r.

43
3-0

I 3-5

\— (b)

"1 Photon energy (eV) —

E.RR. spectrum betore

irradiation

(c)

Figure 3. The resonance absorption at 300 °K due to photoexcited triplet state of the F' -centre in CaO.

Only a single line spectrum is observed since the /W, = 1 levels are not split in cubic

field. The line occurs at g= 1.995 and its presence is accompanied by a strong long- lived fluorescence. The exciting light was obtained from a high-pressure mercury lamp

suitably filtered to cut out all energy below 2.5 eV.

and optical absorption spectrum are readily observable. Although there is a general decrease in the concentration of point defects due to recombination processes, other defects are also produced as a result of thermally induced defect migration.

Wertz et al. [1], have reported the magnetic reso- nance spectra of F2-centres in MgO, CaO, SrO,

and BaO. According to the model proposed by

Professor Wertz and his associates [1], the defect

consists of a single electron trapped at a positivenegative ion vacancy pair. Thus the g-tensor is

expected to be axiaUy symmetric. At X-hand the

and g^i

I -components cannot be resolved with any

great precision: we have, however, resolved them

at ^-band. For MgO our measurements give

> g\\, the respective values being gj^ = 2.0012 and

^1 = 2.0004. For CaO we obtain = 1.995 and

gj^ = 1.9980. The axially symmetric g-tensors substantially support the proposed model. The hyperfine constant for the F2-centre in MgO is somewhat

larger than for the simple F-centre. Surface centres

similarly observed in MgO and CaO powders also have axial symmetry [13]. A comparison of the g- values of the surface centres (see ref. 13) with those

recorded here indicates that the closeness of these centres to the surface merely enhances the crystal

field effects, resulting in larger g'-shifts.

The approximate eigenstates of the F-i-center

have been calculated by Stoneham [14] using a di- electric continuum model in which the electron is bound to an electric dipole field. This is essentially

the approach used in earlier computations by Her- man, WalUs and Milnes [14]. As is common in F-centre calculations, Stoneham uses the high

frequency dielectric constant e since use of the

static constant Co yields no bound states for the

Fa-centre. The principal conclusions are that the

binding energies are 2.42 eV in MgO, 1.92 eV in CaO, 1.70 eV in SrO and 1.24 eV in BaO, and that no other bound states exists within ~ .01 eV of the

conduction bands. The reduced binding energy in the case of CaO and SrO is presumably reflected

in the decreasing importance of surface F-centres

in the overall oxygen absorption processes on sur- faces of these oxides [13]. These calculations,

although only approximate, do give important infor- mation about the optical absorption properties of

F2-centres since it is expected that optical absorption by the defect results in electron excitation into the conduction band. Consequently, it seems probable that the Fa-centres may be optically bleached

using photons of the appropriate energy. King and

Henderson [15] have observed such bleaching in the MgO Fa-centre by monitoring the changes in intensity of the Fa-centre spin resonance signal

during irradiation from a high pressure mercury

lamp. A peak in bleaching efficiency is observed at photon energies of about 3.5 eV as shown in figure 4. Dichvoic effects in bleaching with polarized light were not observed in accord with expectations for an excitation from a bound state to the conduction

band.

3.2. Studies of Zero-Phonon Lines in

Annealed Crystals

The change in optical absorption spectrum which

follows anneahng at temperatures above 300 °C

44
A lOKF SCALE)

BLEACHING EFFICIENCY OF FENERGY (ev) CENTRE IN NEUTRON IRRADIATED MAGNESIUM

OXIDE

Figure 4. The bleaching of the F2-centre in MgO.

Monochromatic radiation was obtained by filtering the radiation from a high pressure mercury lamp. Allowance was made for the strong emission lines in the radiation and for the self absorbance of the specimen. The peak efficiency of bleaching occurs at about 3.7 eV, which corresponds to the separation between the ground state of the /^z-centre and the conduction band. The symbols are: O close to strong emission line in the radiation; A not close to strong emission hne.

was first discussed by Wertz et al. [16]. At low

temperatures the complex absorption spectrum

showing many overlapping bands is highlighted by

a large number of zero-phonon hues (fig. 5). The

technique of uniaxial stress sphtting, discussed

independently by Kaplianskii [19] and by Runci- man [18] has been used to investigate the symmetry properties of defects giving rise to these

lines. Using such techniques it has been shown that zero-phonon hues at 4684 and 4700 A are due to defects having identical symmetry properties [15].

The analysis of the experimentally determined sphtting pattern for these Mnes, shown in figure 6, confirms that the defects have [110] orientation,

orthorhombic symmetry and a Hnear electric dipole

oscillator along a [110] direction. Such symmetry

properties are identical with those expected for M-centres. The stress sphtting patterns of the Unes at 5248,

6419, and 6490 A show that they are due to A

E

transitions within defects having trigonal symmetry

and [111] symmetry axes. Thus the defects have

properties analogous to the /? -centres in the alkali halides [8], and models for these centres are cor4.,000 6.000 8.000 10,000 WAVElEMGTH (&) Figure 5. The optical absorption spectrum of magnesium oxide neutron irradiated to a dose of 2 X 10^" nvt at 600 "C.

This spectrum has many identical features of samples irradiated at 150 °C and

annealed at 300 to 600 °C.

respondingly similar [19]. Figure 7 shows that the

defects consist of three anion vacancies forming an

equilateral triangle on the [111] plane and containing different populations of trapped electrons.

It seems that from electrostatic consideration the

45
P [lOo] p [no] L [no] THEORETICAL EXPERIMENTAL

4660 4690 4700 S. 4680 4690 4700 X

4680 4690 4700 H 4680 4690 4700 I

(<) (") SPLITTING OF 4685X AND 4700X ZERO-PHONON LINES

IN NEUTRON IRRADIATED M90((a) 46 kQi/mmJ (b') 4 1- S kji/xni^)

Figure 6. The experimental splitting patterns for the 4684 and

4700 A lines for two directions of applied stress at 77 °K.

(tt refers to E parallel to stress, tr to £ perpendicular to stress.)

centre will trap more than 3 or 4 electrons; more

likely is the recent suggestions [19] that such a vacancy aggregate will trap 5, 6, or 7 electrons. Using a molecular orbital approach to construct the

most hkely one-electron ground states we obtain:

5 electrons (la)2(le)3

R 6 electrons-M, (lai)2(le)4

R- 7 electrons -2^1 {lai)Hle)^2a,)

.

The hne at 5248 A is apparently due to the 5 elec- tron centre since at low temperatures it is strongly

dichroic under uniaxial stress [19]. This behavior

is associated with the stress induced removal of

the ground state orbital degeneracy, and the in- tensity ratios between the two optical transitions

/ Q- /

5248 A lin« 6419 A line 6490 A line

Figure 7. This shows the proposed models for the three defects whose optical absorption is due to A E transitions within

defects trigonal symmetry.

20 cnT'

1-4

STRESS INDUCED DICHROiSM OF 5248 i ZERO-PHONON LINE IN NEUTRON IRRADIATED M9O ((lOO) STRAIN 50h9S/mm^)

Figure 8. The stress induced dichroism in the 5248 A line under

[100] strain (50 kg/mm).

(a) Dichroism expected for an E—*A transition of a trigonal centre with (110) and

(112) dipoles under [100] stress (b) Experimentally observed behaviour of the 5248 A

Une.

will then be proportional to exp kT

where

AE is the stress induced sphtting in the ground

state. The comparative intensities under stress at 77 and 4 °K are shown in figure 8 and are consistent

with those expected for a trigonal defect with optical dipole moments along [110] and [112] directions in the plane perpendicular to the trigonal axis. No such dichroism is observed for the 6419 and

6490 A lines, and these have been assigned to the

6 electron centre and 7 electron centre, respectively [19, 20].

The description of the electronic ground states discussed above suggests that both the /?+ and R~

centres are paramagnetic. Kemp et al. [21J, have

shown that such is the case for the /?+ centre using

Faraday rotation techniques. The magnetic properties of this centre are expected to be similar to the /?-centre in KCl discussed recently by Krupka

and Silsbee [22]. In fact, strong coupling of the

ground state to the vibrations of the lattice lead

to an unusually short relaxation time in KCl which

explains the absence of the resonance signal even

at low temperatures. The paramagnetic resonance spectrum of the /?-centre in KCl is observed only

under strong uniaxial pressure and at 2 °K. Thus the

previous suggestion [19] by the author that a spec- trum in MgO with trigonally symmetric g- tensor was

due to the R^ centre is likely to be incorrect. However, the 7 electron centre R' which has a Mi ground

state is not expected to present such problems and

therefore it would seem appropriate to assign this spectrum to the /? "-centre. The spectrum shown in figure 9 is recorded in the [110] and [100] orientations and consists of four equally intense lines.

46
The presence of hyperfine interactions with Mg^^

is shown clearly, although it has not yet been fully

investigated. The g'-tensor has axial symmetry, the axes of symmetry being [111] - We obtain g||= 2. 0026

and gj^ = 2.0039. It is difficult to prove that this spec- trum is due to the /?-likc centre containing 7 trapped electrons although the circumstantial evi- dence is quite strong. It is evident from figure 5 which shows the optical absorption spectrum from

the same sample, that the 6419 A hne is much more

intense than the other trigonal centres. Further

studies of this centre are currently in progress.

4. A Trapped Hole Centre in AI2O3

Although optical absorption and spin resonance

studies of irradiation induced defects in alumina

have been reported by numerous authors [3, 23, 24], no correlation has been made of the various optical

bands to the paramagnetic resonance spectra. In reactor irradiated AI2O3, a twelve hne spectrum

from a triplet state in six inequivalent sites was

attributed by Gamble et al. [3], to the molecule ion

AlO^" situated in a pair of oxygen ion sites. The

present results, and those of Cox [25], suggest

that a more probable model, involves a pair of holes

trapped on adjacent oxygen ion sites: the line joining the two 0~ ions defines the paramagnetic z- axis, while the jc-axis points towards a cation va- cancy. It is apparent in figure 10, that the hole pair

may be trapped on any two of three equivalent

anions. Thus there are three equivalent magnetic

2-axes for the defect which are rotated through

+ 3° with respect to the a-directions of the crystal. In other AI2O3 molecular units in the lattice these axes are rotated through — 3° so that in all there are

six inequivalent sites for this defect giving a twelve

line spectrum for a spin triplet state. Clearly this model has the same symmetry as the A10^~ molecule ion discussed by Gamble et al. [3] and the

evidence against the molecule ion model is obtained from temperature dependent properties of

the spectrum.

20023

H JK.

Figure 9. Derivative of the resonance absorption at 35 Gcs ofthe magnesium oxide crystal used to obtain the optical absorption spectrum

shown in figure 8. The spectrum was measured at 77 °K with //„ parallel to (a) a [100| crystal axis and (6) a [110| crystal axis.

47
C - AXIS

0

I

Hie -axis

= 22° / Z-AXIS

Figure 10. Symmetry properties of the hole-pair trapped on

neighbouring oxygen ions in AI2O3.

Figure 11 shows the spectrum measured at 77

and at 300 °K. At 77 °K it is identical with that re- ported earlier [3]. When the Zeeman field is per- pendicular to the crystal c-axis the spectrum consists of 12 equally intense resonance Unes whose

widths were of order 30 to 35 G depending on the

irradiation dose. The spectrum corresponding to one orientation with the field H parallel to the c- axis consists of only two hnes: this spectrum re- mains unaffected when measurements are recorded

at 300 °K. However, the 12 lines of the c-axis spectrum are absent at this temperature. This

effect is concerned with the thermally activated

jumping between the three equivalent hole sites

at the higher temperature. The jump rate increases

with temperature and thus the lines in the c-axis spectrum will broaden with temperature since the

magnetic energy changes when the defect z-axis changes. Clearly this does not affect the H\ \ c-axis

spectrum since in this orientation all possible sites for the defect are degenerate. Cox [25] has noted

that the broadening becomes detectable at 230 °K

and estimates that at this temperature the jump

frequency is of order 10^ sec"'. If we assume weak

dipolar coupling between the hole pair then the

magnetic dipole strength should be approximately

equal to the magnitude of D, the crystal field sphtting. Thus it can be shown that D^U.l^ (7)

J

0^- ion • 0~ ion H II ©

®

Al ion

Al^* vacancy

fl>c =

Figure 11. The derivative absorption spectrum of a reactor irradiated AI2O3 crystal at 77 and 300 °K in two orientations.

Here Be refers to the angle between //„ and the crystal c-axis, and <t>c is the azimuthal

angle measured in the basal plane where 0^=6, = 90° corresponds to the [100] hex- agonal crystal axis. The energy level for the system may be described by a spin- Hamiltonian S^=PH g-S + D{Sl-V3S(S+\))+E(Sl-Sl),

in which S=l, g^ = 2.019, ^-11 = 2.010, 0=710x10-' cm-' and £ = 25X10-' cm"'.

where r is the separation between the holes. We

find since D~ 1400g that r = 3.00 A. This is very

close to the 0-0 separation in the corundum lat- tice, strongly supporting the proposed model for the defect.

It is a pleasure to record my indebtedness to R. D. King of the Ceramic Division, U.K.A.E.A.,

Harwell with whom I collaborated on some of the

work discussed here. I would also thank A. M.

Stoneham, A. E. Hughes, and D. Pooley of U.K.- A.E.A., Harwell, for many stimulating discussions,

and R. Cox of Centre d'Etudes Nucleaires de

Grenoble for communicating his results on AI2O3

to me prior to pubhcation.

5. References

[1] Wertz, J. E., Orton, J. W. and Auzins, P., Disc. Farad. Soc. 30, 40 (1961).

48
[2] Takeda. T., and Watanabe, A., J. Phys. Soc. Jap. 21,

267 (1966).

[3] Gamble, F. I.. Bartram, C. G., Young, C. G., Gilliam, O. R., and Levy, P. W., Phys. Rev. 134, A577 (1964).

[4] Henderson, B., unpublished data.

[5] Weeks, R. A., Interaction of Radiation With Matter, Ed.

A. Bishay, Plenum Press, New York (1967).

[6] Pryce, M. H. L, Phonons (Scottish Universities Summer

School, 1965), Oliver and Boyd (1966).

[7] Klick, C. C., Paterson, D. A., and Knox, R. S., Phys. Rev. 133, A1717 (1964).

[8] Dexter, D. L., JOick, C. C., Russell, G. A., Phys. Rev. 100, 603 (1955).

[9] Hughes, A. E., J. de Phys. 28 (supp), 55 (1967).

[10] Hughes, A. E., private communication.

[11] Ward, W. C. and Hensley, E. B., Bull. Amer. Phys. Soc. 10, 307 (1965).

[12] Kemp, J. C., Ziniker, W. M. and Hensley, E. B., Phys.

Lett. 25A, 43 (1967).

[13] Nelson, R. L., Tench, A. J., and Harmsworth, B. J., Trans.

Far. Soc. 63, 1427 (1966).

[14] Stoneham, A. M., private communication.

[15] King, R. D. and Henderson, B., Pnjc. Brit. Ceram. Soc. 9, 63 (1967).

[16] Wertz, J. E., SaviUe, G., Auzins, P., and Orton, J. W.,

J. Phys. Soc. Jap. Supp. z. 18, 305 (1963).

[17] Kaphanskii, A. A., Optics Spectrosc. 16, 329 and 557 (1964).

[18] Runciman, W. A., Proc. Phys. Soc. 86, 629 (1965).

[19] Henderson, B., and King, R. D., J. de Phys. 28 (supp), 75 (1967).

[20] Ludlow, L K., unpubhshed.

[21] Kemp, J. C, Ziniker, W. M. and Glaze, J. A., Proc. Brit. Ceram. Soc. 9, 109 (1967).

[22] Krupka, D. C. and Silsbee, R. H., Phys. Rev. 152, 816

(1966).

[23] Levy, P. W., Phys. Rev. 123, 1226 (1961).

[24] MitcheU, E. W. J., Rigden, J. D. and Townsend, P. W.,

Phil. Mag. 5, 1013 (1960).

[25] Cox, R., private communication (1967).

49

Covalency in Metal Oxides

B. E. F. Fender and B. C. Tofield

Inorganic Chemistry Laboratory, Oxford University

Recent developments in our understanding of neutron diffraction by antiferromagnetic solids have

made it possible to obtain in a relatively direct way

a quantitative measure of covalency — including the

determination of electron delocahzation in some

simple transition metal oxides. Theory predicts [1] ' that in an octahedral "complex," and for a molecular orbital model involving only d electrons of the

metal and s and p electrons of the Ugand,

S/So = \-a{bAl+ cAl + dAj)

where S is the effective spin quantum number of

the metal ion in an antiferromagnetic solid; So is the expected value of the spin quantum number in a fully ionic lattice; Aa, A^r, and As are covalency or admixture parameters describing cr overlap of d and

p atomic orbitals, tt overlap of d and p orbitals and

(7 overlap of d and s orbitals, respectively; and

a, b, c, d, are constants which depend on the number

of outer d electrons in the cation.

The intensities of neutron Bragg peaks have been

measured to evaluate S and thus the covalency

parameters in MnO [2, 3] and NiO [1, 3, 4] (see

table).

Table. Covalency parameters

NiO (A% + A^,} *6.0% [4] 3.8 ±0.2% [3] MnO (A% + 2A% + Al) *3.3% [2] 3.5 ±0.6% [3] CoO (Al^'^Al + Ai^ 2.8<7o

*Not corrected for zero point spin deviation.

In a similar way the two independent determinations of the effective spin of Co^^ in CoO [5, 6] can

be used to calculate covalency parameters. Using

the average value of S= 1.915 together with an

appropriate value of the spectroscopic sphtting

factor (g= 4.278 from Co2+ in MgO) [7] and allowing

' Figures in brackets indicate the literature references at the end of this paper.

for magnetic disorder at 0 °K [8], the sum of the

covalency parameters is calculated to be 2.8 percent. The similarity of the sum of the covalency

parameters for all three oxides suggests that the

type of bonding is not important and the dominant

factor is merely a reduction of charge at the cation.

The magnitude of the covalency parameters indi- cates that the charge in the region of the cation is reduced by about 0.2e in each oxide, even though

the TT bonding contribution varies from zero in NiO to > 50 percent in MnO.

If the total covalency in metal oxides is governed

mainly by the electron affinities of the metal ions, one would expect the covalency parameter sum for trivalent ions of the first transition series to be

larger, but similar to each other. The present evi- dence does not appear to support this view. The

results of Nathans [2] when corrected for zero point

spin deviation [3, 8] give A^ = 3.5 percent for Cr^+

in LaCrOa and ^2 + 2^2 + ^2 = 73 percent for Fe^+ in LaFeOa. (The transition metal ions are in a near octahedral arrangement of oxygen ions [13]). A similar single set of measurements on the same

systems [9] appear to confirm the results for Cr'^

{A%~2 percent) but indicate that covalency may

be higher for Fe^+ (~ 18 percent).

It is important to note, however, that both

LaFeOs and LaCrOa has distorted perovskite

structures, though neither have been subjected

to a fuU structural analysis. The present cova- lency parameters, and presumably those of Nathans,

have been calculated assuming the atom positions

in both LaFeOs and LaCrOs to be the same as in La(Mno.8Coo.2)03 [10]. The unit cell parameters

[10, 11, 12] indicate that this may be a better

approximation for LaFeOa than for LaCrOa. As

the determination of covalency parameters by

neutron diffraction depends on an accurate knowledge of the nuclear structure factors, if these are

not well known, the covalency parameters ob- tained may be in error. In addition, deviation of

the compounds from stoichiometry can introduce

a further source of uncertainty. Further progress

towards the determination of covalency in trivalent oxides must therefore be Unked with improved

characterisation of the compounds investigated.

51
References

[1] J. Hubbard and W. MarshaU, Proc. Phys. Soc. 86, 561,

(1965).

[2] R. Nathans, G. WiU and D. E. Cox, Proc. Int. Conf. Magnetism, Nottingham (1964).

[3] B. E. F. Fender, A. J. Jacobson and F. A. Wedgwood, J. Chem. Phys. (to be published).

[4] H. A. Alperin, J. Phys. Soc. Japan, Suppl. B III, 17, 12, (1962).

[5] W. L. Roth, Phys. Rev. 110, 1333 (1958).

[6] D. C. Khan, Ph.D. Thesis, Ohio State University (1965).

[7] A. D. Liehr and C. J. Ballhausen, Ann. Phys. (N.Y.) 6,

134 (1959).

[8] T. Ogushi, Phys. Rev. Ill, 1063 (1958).

[9] B. E. F. Fender and B. C. Tofield (unpublished results).

[10] M. A. Gilleo, Acta Cryst. 10, 161 (1957).

[11] S. GeUer and E. A. Wood, Acta Cryst. 9 , 563 (1956).

[12] S. Geller, Acta Cryst. 10, 243 (1957).

[13] P. Coppens and M. Eibschutz, Acta. Cryst. 19 , 524 (1965).

52
Diffusion and Ionic Conductivity: Kinetic Theory

John R. Manning

Institute for Materials Research, National Bureau of Standards, Washington, D.C. 20234

Kinetic expressions for the tracer diffusion coefficient D* of species t and the ionic conductivity

cr, from species i are derived in terms of atom jump frequencies by making suitable modifications in the random walk diffusion equations. Also, D* and cr, are related to each other. The tracer diffusion

coefficient is found from the Pick's Law term in the kinetic equation for the flux Ji, whereas cr, is derived from the drift velocity term in this equation. Nonrandom effects which must be included are those from kinetic driving forces, correlation effects, and the vacancy wind effect. The correlation

factor / is derived from an equation for the effective frequency of independent displacements. It is shown that isotope effect measurements, which measure a quantity /AK, and/or measurements of

(TilD* can help to distinguish experimentally between different diffusion mechanisms. In the case of impurity diffusion, measurements of/and (TilD* yield information about the vacancy jump frequencies

for exchange with solvent atoms in the vicinity of the impurity. A brief description is given of Lidiard's

pair association treatment for vacancy-impurity complexes.

Key Words: Atom transport, diffusion in crystals, diffusion coefficient, diffusion mechanisms, ionic conductivity, jump frequencies, kinetic theory.

1. Introduction

1.1. Kinetic Equations for Diffusion Coefficient and Ionic Conductivity

In the present paper, kinetic expressions in terms

of the atom jump frequencies will be derived for the tracer diffusion coefficient D* of species i and

the ionic conductivity cr, from this species. By use

of these kinetic expressions, the experimental

Df and en can be related to each other. In addition,

information about diffusion mechanisms and about

jump frequencies near an impurity can be obtained.

Diffusion and ionic conductivity both arise from

atom transport. From a kinetic viewpoint, both of

these effects are proportional to the atom (or ion)

jump frequencies. For planar diffusion in the x

direction, the jump frequencies for diffusing atoms

of species i can be defined to be Fr in the + x direction and Tl in the —x direction. Then, Df is proportional to the sum F^ + Fi,, whereas cr; is proportional

to the difference Fft — Fl. The kinetic equation for the flux Ji of species i contains two terms, one

proportional to F^ + Fl and hence to D*, the second

proportional to Fr — Fl and hence to cr,. Thus,

Df and cTi are closely related and can be obtained

from the two terms in the kinetic diffusion equation

for Ji.

1.2. Tracer Profiles

Df and en also can be related to the moments

of the concentration profile Ci{x, t) for diffusion of

a layer of tracer atoms. Here, Df is proportional to

(x^), the mean square displacement of the diffusing

i atoms, whereas at is proportional to the mean

displacement (x).

If a layer of radioactive tracer atoms originally

at plane 0 is allowed to diffuse by means of a random

walk, the atom concentration Ci{x, t) after time

t is given by

Ciix, t) = SoinDftr^l-'expi-x^lWft), (1)

where x is the displacement from plane 0, So is the original concentration per unit area on plane 0,

and Df is the tracer diffusion coefficient. It is as- sumed here that jumps in the +x and —x directions are equally probable so that the mean atom

displacement (x) equals zero. The concentration

profile for eq (1) is shown in figure la. It is a symmetric profile centered on plane 0. The mean

square displacement {x^} of the atoms provides

a measure of the width of the tracer profile. It

is related to the tracer diffusion coefficient by

{x^)^2Dft. (2)

Whenever the probabihty of jump in the -\-x direction is larger than that in the —x direction,

{x) is no longer zero. If each jump is stiU independent of previous jumps,

Ci{x,t) = So{7TD*t)-'l^ exp [-(x - (x))2/4D*r].

(3)

Here Df is the same quantity as that in eq (1).

53
(a) (b)

Figure 1. Concentration profiles for a layer of radioactive tracer atoms originally at plane 0 with p being the probability ofjump to the right and q being the probability ofjump to the left.

(a) Random walk, no kinetic driving force.

(b) With kinetic driving force.

The concentration profile for eq (3) is shown in figure lb. This profile has the same shape and

width as that for the truly random walk of figure la,

but the whole profile is shifted a distance {x).

Instead of there being only one parameter Df to be measured, there now are two parameters

Df and (x).

An electric field acting on charged ions provides

a driving force which makes +x and —x jumps

unequal in number. The resulting (x) is directly

proportional to the field and to cr,. When a constant

electric field is the only driving force, the ionic conductivity o", for species i can be calculated di- rectly from (x).

1.3. Nonrandom EflFects

The basic equations which will be used in the

following discussion are the random walk dif- fusion equations. The introduction of a driving

force by itself introduces a nonrandom effect into the kinetic equations. This is easily included in random walk equations by allowing Tr to be un- equal to Vl, as in eq (3) and section 2.1. Even in the absence of driving forces, however, atoms often

do not follow random walks. Instead, if diffusion occurs by a vacancy or interstitialcy mechanism,

they follow correlated walks, where the direction

of one atom jump depends to a certain extent on

the directions of previous jumps by that atom. This

introduces a correlation factor / into the diffusion

equations and influences both Df and (x). In addition, if there is a net flux of vacancies or interstitialcies, the influence of the flux will change the

atom jump frequencies. This can greatly affect (x) but does not influence Df.

The modifications of the random walk equations

introduced by correlation effects and a vacancy

flux are considered in separate sections on the

correlation factor and the "vacancy wind". These

nonrandom effects make the relations between

the experimental quantities (Df and cr,) and the

atom jump frequencies somewhat more complex.

Nevertheless, they also can help one to distinguish

experimentally between different diffusion mechanisms and to calculate jump frequency ratios for jumps in the vicinity of an impurity.

2. Kinetic Diffusion Equations Assuming Independent Displacements

2.1. Random Walk Diifusion Equations

Modified by a Driving Force

Although atoms do not necessarily follow random

walks in actual crystals, many of the random walk

ideas can be apphed to correlated walks. As a first step, the random walk diffusion equations will be

discussed.

The random walk diffusion equations modified

by a driving force can be derived by considering

two neighboring atom planes which are normal

to the diffusion direction. The number of atoms of

the diffusing species i per unit area on plane 1 can

be called ni and that on plane 2, which is the

neighboring plane in the + x direction, can be called

ri y. Then, if is the jump frequency for a jump of

a given i atom from plane 1 to plane 2 and is the jump frequency in the opposite direction from

plane 2 to plane 1, the net flux Ji of species i in the + x direction is

y, = niri2 — n2r2i. (4)

Here niTvi is the number of atoms jumping in the

+ x direction per unit area per unit time and

n2r2i is the number jumping in the —x direction. Equation (4) can be written in terms of two

differences: (1) the difference between the concentrations Hi and R2 and (2) the difference between

the jump frequencies ri2 and r2i. If planes 1 and 2

54
(and all such neighboring planes) are separated

by a distance d

n = cd, (5)

where c is the number of atoms per unit volume.

Also the difference n2-~ni equals d times the

gradient driildx. Thus,

n-i — n\ = d^{dcildx) . (6)

Equation (4) then reduces to

Ji = -TdHdCildx) + Cid(Yri -T-n), (7)

where Y is the average of and r2i, equal to 2(ri2 + r2i), and c, is the average concentration

i d{ni -\- ni) , equal to the concentration associated

with the region midway between planes 1 and 2. A more famihar form of the diffusion equation can be obtained by writing

Df^Yd\ (8)

and »

{vF)i = d{Yn-Y.n). (9)

Here, D* is the tracer diffusion coefficient for species i measured in a homogeneous system, and

{vF)i is the average drift velocity of species i as

a result of kinetic driving forces acting on the atoms. (By definition, a kinetic driving force is any

influence which makes the jump frequency ri2 in one direction between two planes differ from the

frequency r2i for a jump in the opposite direction

between these same two planes.) This yields

Ji = -Df{dcildx)+Ci{vF}i. (10)

The first term on the right in eq (10) is the normal

Pick's Law diffusion term with Df being related to the width of the tracer profiles shown in figures la

and lb. The second term is the drift term with

{vF)i being related to the displacement (x) of the

tracer profile shown in figure lb and also related

to the ionic conductivity o"; for this species.

2.2. Effect of Diffusion Coefficient Gradient

In the derivation of eq (10), it was assumed that each atom jump was independent of previous

jumps; but no assumption was made about the

relation between r23 and ri2, where r23 is the jump

frequency in the +x direction from plane 2 and

ri2 the similar frequency from plane 1. If the jump

' Figures in braclcets indicate the literature references at the end of this paper.

frequency depends on the starting position as well

as the direction of the jump, r23 and Y12 may differ. This circumstance can arise, for example, when

the energy of motion (and/or formation) depends

on the composition of the crystal and the composition is a function of position. This would make the

diffusion coefficient depend on position, and conceivably could strongly affect the diffusion process. Nevertheless, it can be shown that such a depend- ence of diffusion coefficient on position does not

affect eq (10). Thus, this equation is vahd in general, even when there is a non-constant diffusion

coefficient. One quantity which is affected by a diffusion

coefficient gradient is the total drift velocity {v)i

of the i atoms. This velocity depends on differences,

such as r23 — r2i, between the frequencies for jumps in the +x and —x directions from a given

plane (in this case, plane 2). If ri2 ^ r23, the ex- pression for {v)i differs from the expression for (uf);. In general, [1, 2]'

{v)i^ {vF)i + {dD*ldx), (11)

where dD*ldx is the diffusion coefficient gradient.

The measured drift velocity {v)i equals the drift velocity (ff ); arising from the driving forces

only if dDfjdx equals zero. Df is defined by eq (8), which is vahd whenever

individual displacements are independent of one

another. In general, Y and hence Df can depend

on local conditions in the crystal. Nevertheless, r and Df often do not depend on position, making

dDfjdx equal zero. For simplicity, it is assumed in the following two sections (3 and 4) that Y is independent of position.

3. Correlation Effects

For diffusion by a vacancy or interstitialcy mechanism, neither eq (8) for Df nor eq (9) for {vF)i is accurate if F is assumed to equal the actual jump

frequency. The correlation factor influences

Df, and both the correlation factor and vacancy

wind effects influence (vf)}- In the present section

correlation factor effects are considered.

3.1. Correlation Factor

After an atom has exchanged places with a given

vacancy, the vacancy is in the proper position to allow the atom to make a reverse jump back to

its original site. Thus, the second exchange of an atom with a given vacancy is not random in direction. Instead, the atom has a greater-than-random

probabihty of making a reverse jump. This situation, which arises whenever diffusion proceeds by

55
a vacancy mechanism, creates a correlation between the directions of successive atom jumps

and introduces a correlation factor into the dif- fusion equations.

A similar correlation arises for diffusion by an

interstitialcy mechanism. After an atom has jumped

from an interstitial to a lattice site, the atom which was just displaced from the lattice site is still on a neighboring site and is in the proper position to cause a reverse jump. By contrast, no such correlation arises for self-diffusion by an exchange mechanism or by a direct interstitial mechanism.

Mathematically, the correlation factor can be cal- culated as follows: In general, regardless of whether

there is a random walk, eq (2) is vahd in the hmit

where t goes to zero [3]. The total displacement x

of an atom is the sum of the displacements xj from

the individual jumps made by the atom,

X — Xi+ X2 + X3+ . . . + Xn- (12)

Thus,

x'=J^xj + 2''2 i xjx,. (13)

If each jump is independent of previous jumps,

the summations of xjXk on the average equal zero. For this case,

n ^random = (20"' ^ xj. (14)

J = l

In cubic crystals, eq (14) reduces to

Z)F^"'*°'" = ^ (15)

where k is the jump distance and v is the atom

jump frequency (including jumps in all possible

directions).

For a correlated walk, one obtains instead

£)* =yZ)random_

Here the correlation factor / is defined as equalling

the ratio Df/Dp"^"'"; and, from eq (13),

II -I n ^^'p"' (17)

In cubic crystals, it follows that the tracer diffusion

coefficient can be expressed simply as

Dt^lk^vf. (18) .

Equation (17) is the equation which is usually | used to calculate the correlation factor. A number j of investigators have calculated values of / for ! various circumstances from this equation [4]. In particular, very general matrix equations have ' been developed by Howard [5] . i

3.2. Correlation Effects in Terms of Effective ; Jump Frequencies

The calculation of /from eq (17) is directly appH- ' cable only to the Df or (x^) terms in the diffusion

equations [i.e., the width of the tracer profile in ; figure 1 and first term on the right in eq (10)]. By

itself, eq (17) does not provide information about : the effect of a correlated walk on (vf); or (r) [i.e., I the shift of the tracer profile in figure lb and the ' second term on the right in eq (10)]. In addition, \ the mathematical equations leading to eq (17) do

not provide a convenient physical model for dis- cussion of correlation effects. For these reasons, an alternate but still rigorous approach for calculation of correlation effects will be presented here. The majority of the discussion will be in terms of j diffusion by a vacancy mechanism, but the same

reasoning can be applied to diffusion by an inter- ; stitialcy mechanism. ; If diffusion normal to a set of sufficiently close- , packed planes is considered, the correlated walk

equations can be expressed in terms of an equiva- \ lent random walk with the actual jump frequencies, ; and Foi, in eqs (8) and (9) being replaced by i effective frequencies of independent displacement,

Ffj and r.f, . Each close-packed plane must be ' such that a vacancy on one side of the plane cannot

pass to the other side of the plane without at some

'

time stopping at a lattice site on this plane. In cubic crystals, the (100) planes are suitable planes.

Any vacancy which arrives on the same plane as ; a particular tracer atom can be said to have randomized its position with respect to the tracer, since ! in the absence of driving forces the vacancy is just as likely to cause future jumps of the tracer ! in the —x direction as the + x direction. Even with ! a driving force, the distortions of the vacancy flux

j introduced by the tracer will not change the vacancy , concentration on this plane. Thus, vacancy paths j which start from this plane can be treated as being i uncorrelated to previous jumps of the vacancy and

the tracer. With this background, let us consider a series of i exchanges of a tracer with a given vacancy V. ! After an initial exchange which moves the atom a distance d in the +x direction, the vacancy V is i
in the proper position to allow a reverse jump

taking the atom back to its original plane. There

is probabiHty P that such a jump will be caused by

vacancy V (this may occur either immediately or

after the vacancy has wandered through the lat- tice for some time) and probabihty 1— F that the

vacancy instead moves to a vacancy sink or to a

site on the plane containing the tracer. In these

latter cases, the correlated series of atom-vacancy

exchanges is terminated and the next tracer jump

starts a new jump series which is uncorrelated to previous jumps. Unless at some time the corre- lated jump series is terminated in this way, the atom

I and vacancy will continue to jump back and forth between the two neighboring planes, contributing

many atom jumps but resulting in almost no net

I X displacement of the atom.

If a series of correlated jumps contains an odd

number of jumps, there is net displacement d, and

if the series contains an even number of jumps

there is zero net x displacement. The probabiHty

L that a correlated series results in a net x displacement of magnitude d is given by

U=l-P^P-^-P^+ . . .^i\ + P)-\ (19)

Here the nth term in the series equals the probabihty that the correlated series contains at least n jumps.

Only a fraction of the jumps wiU be the first jump

in a series. This fraction equals \ — P, the probabihty that another series terminated on the pre- vious jump. Therefore, the probabihty that a given

jump starts a series which results in net displacement d equals U{\ — P). Displacements resulting

from separate series are independent of one another. As a result, when there are no kinetic driving

forces, an effective frequency of independent

displacements (all of magnitude d) can be defined

in terms of the actual jump frequency Y" as

p=r«[f/(i-P)] = r'^- (20)

or

Y'=Y"f, (21)

where the correlation factor /is given by

1-F

/=Y^- (22)

For crystals to which the present treatment ap- phes, eq (22) for the correlation factor is the same

as that obtained from eq (17). It can be seen that

the correlation factor equals the fraction of jumps

which result in independent displacements, with

the displacement neither cancehng previous jumps

nor being cancelled by succeeding, correlated

jumps.

Substituting the value of Y^ from eq (21) in place

of r in eq (8) yields

Df = Y"fd\ (23)

For cubic crystals, this is equivalent to eq (18).

3.3. Values of the Correlation Factor

The probabihty P can be calculated from the

crystal geometry in terms of the vacancy jump

frequencies in the crystal. For self-diffusion, P

and hence the correlation factor / are numerical

constants depending on the crystal structure. To

first order, P for self-diffusion equals where z is the number of nearest neighbors to which a vacancy

can jump. More careful calculations show that P

for self-diffusion is somewhat larger than z"^ since even those vacancy paths which do not cause an immediate vacancy-tracer exchange can eventually

allow such an exchange to occur. To two significant

figures, / is 0.78 for self-diffusion in face-centered

cubic crystals (or NaCl sub-lattices), 0.73 in bodycentered cubic crystals, 0.65 in simple cubic crystals (or CsCl sublattices), and 0.50 in diamond [6]. Values of / when diffusion is by an interstitialcy

mechanism differ from those for a vacancy mechanism. For example, in the face-centered-cubic

case,/ equals ^ when diffusion is by a coUinear interstitialcy mechanism. For self-diffusion by ex- change or direct interstitial mechanisms, / equals

unity. This dependence of/ on the diffusion mechanism can help in experimental determination of the

diffusion mechanism.

For anisotropic diffusion, the correlation factor

usually will be different for diffusion along each

principal axis. Values for the aluminum sublattice

in corundum are given in table 1. Here the correlation factors fx and fz depend strongly on the jump

frequency ratio Vzjvx- If/x or fz were measured, the

corresponding value of Vzlvx could be found.

For diffusion of impurities, the correlation factor depends on the vacancy jump frequencies near the

impurity. For example, in a face-centered cubic

crystal (or NaCl sublattice), the vacancy jump fre- quencies can be defined in a 5-frequency model as w\ for a jump which moves the vacancy from one

site neighboring on the impurity to another such

site, Wi for exchange with the impurity, wz for dis- sociation from the impurity by a jump from a nearest neighbor site to a nonnearest neighbor site, for

290-621 0-68—

5

57
Table 1. Correlation factors for diffusion via vacancies in the aluminum sublattice of corundum ^

0 0.33 1.00 10 0.50 0.65

00 0.66 0.36

^ After Compaan and Haven. These and additional values are given in reference [6].

^ Vz = jump frequency along z-direction.

i'j- = jump frequency in xy plane. '^fx=fy = correlation factor for jumps in xy plane,

•/z = correlation factor for jumps along the z-direction.

the reverse of a Ws jump, and wo for all other jumps.

Then

W2

(24)

where F depends on the ratio Wilwo and ranges

from 2/7 to unity [7] , and

2wi + IFwz

2w2 + 2wi + IFw^i

(25)

For a fast-diffusing impurity, where W2 is much

larger than Wi or wa, f can become nearly equal to zero, yielding a large correlation effect. Then the

tracer diffusion coefficient D* in eq (18) is proportional to w\ + 3.5Fw3 rather than to W2 {v itself is proportional to W2, but vfin this case is proportional

to W1 + 3.5FW3). On the other hand, if W2 is much

smaller than w\ or W3, f can become almost unity,

and there is only a smaU correlation effect. The introduction of the additional jump fre- quencies Wi, W3, etc.,. into the relation between

D* and v prevents a direct determination of v from measurements of Df. If the factor / can be

measured, however, not only are calculations of v again possible but also eq (25) provides a relation

between the three ratios W\lw2, w\lwz, and wJwq

(for jumps of solvent atoms near the impurity).

Other relatiqns between these ratios can be found

by measuring ynlDf, as described later in this paper; by measuring D*ID*, where Df is the solvent

diffusion coefficient; or by measuring the coefficient b which describes the change in solvent diffusion

coefficient upon the addition of impurities [8].

Where vacancy-impurity complexes are formed,

dielectric and mechanical relaxation experiments

also can yield relations between these jump fre- quency ratios.

3.4. Isotope Effect

One means of measuring the correlation factor

experimentally is by isotope effect experiments.

Two isotopes of the same diffusing element wiU

have different masses and hence can be expected

to have different jump frequencies. In crystals

where each jump vector lies along an axis of at least two-fold or three-fold rotational symmetry [9]

,

AD"

V)8 W2a

(26)

Here AD*^D^ — D* and Aw2 = W2p — W2a, where

subscripts a and /3 refer to the two isotopes and

W2 is the isotope jump frequency.

According to classical rate theory, W2 is proportional to the inverse of the square root of the isotope mass, making

1. (27)

In applying eq (27) to real cyrstals, however, the masses ma and should be reduced masses which

take into account the masses of the other atoms in the crystal which must move to allow the diffusing

isotopes to jump [10]. The usual procedure is to treat rua and mis as being the actual masses of the

isotopes but then to multiply the right-hand side

of eq (27) by a factor AK to correct for the reduced masses [11, 12]. Thus,

AW

Da =f,AK

1/2 —I -1 (28)

It is predicted [11, 12] that AK will equal approximately 0.5 for diffusion by an interstitialcy mechanism, since approximately half the kinetic energy

for the interstitialcy jump is carried by each of the

two atoms cooperating in the jump. Also, in general, if n atoms participate equally in a jump, such

as by a ring mechanism, AA^ would equal n~^. For

direct interstitial and vacancy mechanisms, value

of AK near unity might be expected since only one atom really makes a jump.

The product f^AK can be determined by measuring AD*IDa with the true isotope masses ma and

m^ in eq (28) being known. For diffusion of small

interstitial atoms, which certainly diffuse by a

direct interstitial mechanism, it is usually found

that ffiAK equals unity within experimental error. Since for the direct interstitial mechanism is unity, it follows that AK also is near unity, as expected. For a vacancy mechanism in a face- centered cubic crystal or NaCl sub-lattice,

equals 0.78. Values somewhat smaller than this

58
have been measured for f/s^K (— 0.7) in NaCl [13],

indicating that — 0.9 in this case if it is assumed

that diffusion occurs entirely by means of free

1 vacancies.

Experimentally determined values of the product

ffi^K can be used to distinguish between different possible diffusion mechanism, since predicted

values of this product for self-diffusion depend

strongly on the diffusion mechanism [14]. Here the

factor AA^, although it makes the relation between

AD*ID* and mplnia more complex, greatly aids in

I distinguishing between the possible mechanism.

* For example, the correlation factors for selfI diffusion by direct interstitial and ring mechanisms

are both unity but the values of AK differ greatly.

For self-diffusion in an isotropic crystal, the

I correlation factors are known numerical constants. Thus, if the diffusion mechanism is known, AK

can be calculated from measured values of fpAK.

For impurity diffusion, however, values of are normally not known. Since only the product of

fpAK can be measured, an estimate must be made

j of AK, perhaps by assuming that it is the same as

j for self-diffusion, if one wishes to determine

from isotope effect measurements.

4. The Drift Term and Ionic

Conductivity

To this point, the discussion of non-random effects has been concerned mainly with the effect of cor- relation on the tracer diffusion coefficient D*, which

appears in the first term on the right in eq (10). Equally important are effects on (ff);, the quantity

which appears in the second term on the right in eq (10).

4.1. Correlation EflFects on (ff);

In the discussion of correlation factor effects in terms of the probabiUty P, it was shown that in the absence of driving forces the actual jump fre- quency F" for a given type of jump is related to an

effective frequency F*" of independent displace- ments. It can be shown [2] that this type of dis- cussion is vahd even when there is a driving force. Thus, one can write

rf2 =n (29)

ri, =fn, (30)

where / is given by eq (22) with P evaluated in the

absence of driving forces. The frequencies F*" of independent displacement can be substituted into eq (9) to yield (ff);. The basic jump frequencies

F* equal the actual jump frequencies F" in the absence of kinetic driving forces multiplied by

[1+ {FiXil2kT)], where F, is the net kinetic driv- ing force, Xi is the x displacement resulting from the

jump, k is Boltzmann's constant, and T is the absolute temperature. These relations can be substituted

into eq (9) with F^ taking the place of F to yield,

W)i=r"fdHFiikT), (31)

where d is the interplanar spacing.

It may be noted that the same correlation factor

appears in eq (31) as in eq (23). Thus, eq (31) can be

rewritten as

{vF)i = -j^, (32)

where D* is the tracer diffusion coefficient measured

in a homogeneous system in the absence of kinetic

driving forces.

4.2. Vacancy Wind Effect in an Electric Field

By definition, the kinetic driving forces are any

influences which make a jump by an atom in one

direction between two neighboring sites more likely

than a jump in the opposite direction between these same two sites. An electric field E acting on the

charged ions in a crystal can provide such a driving

force Ff with

Ff = QiE, (33)

where qi is the charge of the diffusing ion. In kinetic

terms, such a field changes the effective energy of

motion for jumps with and opposed to the field. In addition to the direct force exerted on the

ions, there also can be an indirect force. This

indirect force arises if the field also acts on the

other ions in the sublattice in which a given ion is diffusing. If diffusion is by a vacancy mechanism

and the field drives these other ions in the +x

direction, there will be an equal flow of vacancies

in the —x direction. As a result, vacancies will approach any given ion more frequently from the

+ X direction than the — x direction. This causes the

given ion to begin a series of vacancy-ion inter- changes more frequently with a+x jump than with

a —X jump.

In the discussion of correlation effects, it was

found that a correlated series can be treated as giving either zero net x displacement (if there are an even number of jumps in the series) or a dis- placement equal to that from the first jump (if there

are an odd number of jumps in the series). Detailed

analysis [2] shows that if the first exchange is more

Ukely to be in the + x direction an additional kinetic

59
driving force FP in effect is exerted on the ion, with

the force for an isolated impurity being given by

IJ^iE = (vF)i (36)

Ff = 2qsE{njp), (34)

where Qs is the charge of the other (solvent) ions in the sublattice and (rijp) is a complex kinetic quantity, expressions for which are given in reference

[15]. In kinetic terms, {njp)d equals the average

net X displacement a vacancy which starts its path on a site on the -\- x side of the impurity would

undergo upon traveling to a vacancy sink if ex- changes with the impurity were not allowed and

there were no kinetic driving forces. This same

equation applies for diffusion by an interstitialcy

mechanism if "interstitialcy" is substituted for "vacancy" in the definition of (rijp). If the only driving forces are those arising from

the electric field, eqs (32) to (34) yield for a vacancy

mechanism

ED*

(35)

Here the first term inside the parenthesis arises

from the direct force on the charged ions, and the

second term from the influence of the vacancy

flow. This latter effect can also be called a "vacancy

wind" effect. Vacancy wind effects arise whenever there is a

flow of vacancies. The relation of vacancy wind

effects to correlation effects is discussed in refer- ence [16]. For self-diffusion, equals Qs and

1 + 2 (njp) =/"', where /is the correlation factor. In eq (35) correlation effects introduce a factor /

into the expression for Df, whereas the vacancy

wind effect makes the factor in parentheses equal

f~^. Thus, in the expression for {vF)i, the factor

arising from vacancy wind effects exactly balances

/ in D* arising from correlation effects. In other

words, the vacancy wind and correlation factor

effects on (ff); exactly cancel one another for self- diffusion, being equal and opposite effects in this case. For impurity diffusion, however, these effects usually do not cancel each other. For impurity dif- fusion, the vacancy wind term '2qs{njp) makes the

factor in parentheses in eq (35) differ from qi but

usually does not make this factor equal exactly

qif~^. Explicit equations for impurity diffusion are given in section 4.4.

When there is a constant electric field and no driv- ing forces except those from the electric field, the

shift (x) in the tracer profile shown in figure lb

is given by

{x) = fJLiEt. (37)

Also, the ionic conductivity at contributed by

species i is

ai = Ciixiqi, (38)

where c; is the number of i atoms per unit volume

and qi is their charge.

Experimentally, yu,, could be determined from

either eq (37) or eq (38). The total conductivity tr

is given by

(39)

where the summation is over all charge carriers

in the crystal, including electrons. Thus, the frac- tion of the conductivity contributed by species i must be known before measurements of or can be

used to determine /a;. For diffusion by a vacancy mechanism, eqs (35)

and (36) yield

Df kT

l + 2(n,-,)^

Qi- (40)

The similar equation for diffusion by a coUinear

interstitialcy mechanism (where both cooperating

atoms move along the same line) is

Mi + qs Dt kT (41)

Here the direct force affecting a jump of the impurity is E(qi + qs), since a solvent ion must also move during the interstitialcy jump.

In the special case of self-diffusion, the quantities inside the square brackets in eqs (40) and (41) equal where / is the correlation factor [15]. Then, for a vacancy mechanism,

fJ-i ^ Qi Df kTf

(42)

4.3. Drift Mobility and Ionic Conductivity

The ionic drift mobility /x; of species i in an electric field is defined as the average drift velocity

{vi.-)i per unit field. Thus,

For self-diffusion (where qi — qs) hy a collinear interstitialcy mechanism,

fjLi _ 2qi

Df^Wf ^^^^

60
Equation (43) for the collinear interstitialcy mechanism differs by a factor 2 from eq (42) for the va- cancy mechanism. In addition, the numerical values

of the correlation factor depend on the diffusion

mechanism, being 0.78 for a vacancy mechanism

2

in a face-centered cubic structure and - — 0.67

for a coUinear interstitialcy mechanism in this structure. Thus, measurements of fiilDf provide

information about the diffusion mechanism.

For self-diffusion, the ratio iXijDf can be derived

by noting that the net flux of ionic charge in a given

sublattice is proportional to the vacancy or interstitialcy flux. When there are no impurities, motion

of these defects is not correlated; whereas motion

of the ions, which contributes to Df, is correlated. Thus, the ratio of fJii/D* contains the factor /"'. In

addition, the jump vector of the interstitialcy imperfection differs from that of either ion cooperating

in the jump. Thus, the ratio iXilD* also contains a displacement factor S, where

urement of Sf'^ can help to estabhsh the diffusion

mechanism.

As in the case of the isotope effect, it is not just

the correlation factors which enter here. Instead

there is an additional factor, S in this case and A/^

in the isotope effect case, which has an important

influence on the results.

Table 2. Correlation and displacement factors for self-diffusion in silver chloride

Diffusion mechanism S =S/-'

Direct interstitial 1.00 1.00 1.00

1.28 1.00 1.28 Noncollinear interstitialcy

(6 s 70.5°)" 1.03 1.33 1.37 Collinear interstitialcy (^=0)" 1.50 2.00 3.00

angle between the jump vectors of the two atoms co- operating in the interstitialcy jump.

^ _ ^defect^defect'^defect ^atom^atomf^atom (44)

Here \ is the jump distance of the atom or the defect

during an elementary jump, v the jump frequency

per atom or defect, and c the number of atoms or defects per unit volume. For interstitialcies, there

are two atom jumps for each interstitialcy jump,

making i^atomCatom = 2i'defectCdefecf Also, for a collinear

interstitialcy mechanism, ^defect — 2Xatom. making

S = 2. On the other hand, for a vacancy mechanism, S= 1. In general, equations such as (42) and (43) relat- ing fjLi to D* for self-diffusion by means of isolated

defects can be written as

fM ^ qjS

Df~kTf (45)

Also, one can define a quantity Da^ in terms of the

measured conductivity cr,

crkT

CiQi (46)

If only one species contributes to the ionic conductivity and eq (45) is vahd for this species,

Df Sf-K (47)

Values of S, and Sf~^ for various mechanisms

are hsted in table 2. Because of the difference in these values for the different mechanisms, a measMeasurements on silver diffusion in AgCl and

AgBr indicate that Sf~^ is less than 3.00 but larger

than 1.37. It is thought that a combination of col- hnear and noncolhnear interstitialcy jumps is present in these cases [17].

If diffusion proceeds in part by means of bound

anion vacancy — cation vacancy pairs, the motion

of these pairs will contribute to Df but not to fXi or Da- This situation would make (jliID* and DJDf

smaller than the values predicted in eqs (45) and

(47). Thus, if vacancies are known to be the dominant defect, values of DJDf (or fx.,kTjDfqi) smaller

than the value for a vacancy mechanism may

suggest the presence of vacancy pairs, as for example in TlCl [17].

On the other hand, larger values of D^/Df may

be obtained if the total conductivity cr is larger than

CTj, as will occur when electrons or ions on the other

sublattice carry appreciable current. In such a case,

it is necessary to know the transport number of

species i and substitute cr; in place of cr in eq (46)

if one wishes to measure Sf~^. The ratio jXilDf

may provide a more definite measurement of S/~'

in this case, since it is not affected by these other

contributes to cr.

4.4. Drift Mobility for Impurity Diffusion

For diffusion of impurities, the quantities in square brackets in eqs (40) and (41) in general do

not equal and can differ greatly from this value. For example, in a NaCl sublattice, the vacancy jump

frequencies near an impurity are defined above

eq (24). For these frequencies [15],

61
4

(a) (b) (c)

Figure 2.

(a) Value of IF, which appears in eqs (24) and (48) to (52), as a function of wjwo and waluit.

(b) Value of 01. which appears in eq (49), as a function o{ w^iwa and WoIwa-

(c) Value of (ft2, which appears in eq (51), as a function of w^liva and uidIw^.

Wo

— 2wi + WslS — l {1— F) {W4 — Wo)W4 ^] 2wi + 7Fw3

(48)

For a divalent impurity diffusing in a divalent

cation sublattice, qi = qs = 2e. With diffusion by a vacancy mechanism on a sublattice of a crystal with

the NaCl structure, eqs (40) and (48) then yield

rhere

jjj__2e —2wi + (l)iiV3 Dt~kfl2wi + lFw3 .

0, = 6 + 7F - 14( 1 - F) (m;4 - wo)w4'.

(49)

(50)

Values of 7F and which depend only on the

ratio Wilwo, are presented in figures 2a and 2b. For a tetravalent impurity diffusing in a divalent

cation sublattice, g, = 4e and qs = 2e. With diffusion

by a vacancy mechanism on a sublattice of a crystal

with the NaCl structure, it follows in this case from eqs (40) and (48) that

jM

D1

4e

kT

(f)2Ul3 2w\ + lFw3

/here

<^2 = 3 + 7F — 7( 1 — F) {wa — ivn)wi

(51)

(52)

Values of 7F and <^)2, which depend only on the

ratio WaIwq, are plotted in figures 2a and 2c. Here

it is assumed that there is a uniform vacancy concentration much larger than the concentration of

tetravalent impurities i.

4.5. Cross Terms in the Thermodynamic

Diffusion Equations

It can be seen from eqs (48) to (52) and figure 2

that the vacancy wind contribution to fXi/Df [the

term containing {nj,j) in eq (40)] can be large and

in fact can even be larger than the contribution

from the direct force exerted by the electric field on the charged ions. As a result, the deviation from

the Nernst-Einstein relation {fiilDf = qi/kT) can be

large.

For a system with two atom species {A and B)

and with vacancies in equilibrium, the general

thermodynamic diffusion equations can be written

as

JA — LaaXa + LabXb ,

Jb — LbaXa + LbbXb (53)

Where Xa and Xb are the thermodynamic driving

forces for species A and B. Howard and Lidiard

[18] have shown that, if the cross-coefficient Lba

is non-zero, deviations from the Nernst-Einstein

relation can be expected. In the present equations,

the quantity 2{njp) in eqs (40) and (41) equals

LbaILbb, and Lbb for diffusion of the dilute impurity

B equals D%CBlkT. Thus, explicit values of Lba

can be determined.

The vacancy wind contribution to (yf); appears

here entirely as a contribution to Lba- As expected,

this term causes deviations from the NernstEinstein relation. Even for self-diffusion, this cross term is not negligible, since then 2{njp) equals

(1 -/)//•

62
5. Vacancy-Impurity Complexes

When an aliovalent impurity is introduced into an ionic crystal, vacancies may also be introduced

in order to maintain charge neutrality, making the

concentration of vacancies depend on the concentration of impurities. In this case, it is convenient to consider the vacancy-impurity complexes as the

diffusing units, with the probability that an impurity

is part of a complex being a function of impurity

concentration. Lidiard [19] has treated in detail the diffusion of divalent impurities in monovalent

sub-lattices (which is similar to tetravalent impurities in divalent sub-lattices). Two effects which

enter these calculations are the diffusion potential

and the variation with impurity concentration of

the probability p that an impurity is associated

with a vacancy in a complex.

The diffusion potential arises because vacancies

tend to diffuse down their gradient faster than do

the impurities, causing a separation of charge and

hence a local electric field. The resulting diffusion

potential prevents a further separation of charge.

This keeps the vacancies in the vicinity of the

impurities and thus allows the diffusion problem

to be treated entirely in terms of the vacancyimpurity complexes. The homogeneous thermoelectric power arises from the same type of internal

electric field as does the diffusion potential. In the case of the homogeneous thermoelectric power,

there is a separation of charge by diffusion of two

kinds of defects, cation and anion vacancies if there is Schottky disorder or vacancies and interstitials if there is Frenkel disorder.

The variation of p with concentration can have a strong effect on the diffusion coefficient. One can

distinguish two types of experiments: (1) Tracer

experiments can be performed where the number

of aliovalent ions and accompanying vacancies

introduced with the tracer is small compared to the vacancy concentration already present in the

crystal. This already-present vacancy concentration

may be the intrinsic concentration for a given temperature; or, in doped crystals, it can arise from

vacancies introduced by other aliovalent ions

present in uniform concentration throughout the

crystal. In the latter case, p depends on the uniform

concentration of aliovalent ions and not on the

tracer concentration. Since the jump frequency

V equals W2P, where wz is the vacancy-impurity

exchange frequency in the complex, eq (18) becomes

Df = -^K-'w2pf.

in which a large number of ahovalent ions are intro- duced along with the radioactive tracer ions as

carrier (being nonradioactive ions of the same

chemical element as the tracer ions), making the

additional vacancy concentration much larger than

that normally present. In this Hmit, Lidiard [19] has shown for a divalent impurity diffusing in a monovalent sublattice that

ji=-Dr

2 da

1+ p dx (56)

(55)

(2) In the other Umit, experiments can be performed

The factor 2/(1 +p) arises because the concentration gradient of complexes is not directly proportional to the concentration gradient of tracer. Also, p is determined by the total concentration

of aliovalent ions, i.e., tracer plus carrier, in the

particular part of the crystal.

Lidiard's pair-association model [19] is especially

useful in treating vacancy-impurity complexes

and in deriving equations such as eq (56) above.

Also, this model can be applied in general to dilute alloys if dissociation of vacancy-impurity

complexes is taken into account. When this is done,

the pair association model and the correlated walk

model presented earlier in this paper give identi- cal results.

6. References

[1] A. D. LeClaire, Phil. Mag. 3, 921 (1958).

[2] J. R. Manning, Phys. Rev. 139, A126 (1965).

[3] A. Einstein, Ann. Physik 17, 549 (1905).

[4] See for example discussions in Y. Adda and J. Philibert, La Diffusion dans les Solides. (Presses Universitaires de France, Paris, 1966) Chapter 7; or in J. R. Manning,

Diffusion Kinetics for Atoms in Crystals, Ch. 3 (D. Van

Nostrand Co., Inc., Princeton, 1968).

[5] R. E. Howard, Phys. Rev. 144 , 650 (1966).

[6] K. Compaan and Y. Haven, Trans. Faraday Soc. 52, 786

(1956); 54, 1498 (1958).

[7] J. R. Manning, Phys. Rev. 136, A1758 (1964).

[8] R. E. Howard and J. R. Manning, Phys. Rev. 154 , 561

(1967).

[9] A. H. Schoen, Phys. Rev. Letters 1, 138 (1958); K. Thar- malingam and A. B. Lidiard, Phil. Mag. 4, 899 (1959).

[lOJ G. H. Vineyard, J. Phys. Chem. Solids 3, 121 (1957).

[11] J. G. MuUen, Phys. Rev. 121, 1649 (1961).

[12] A. D. LeClaire, Phil. Mag. 14, 1271 (1966).

[13] L. W. Barr and A. D. LeClaire, Proc. Brit. Ceram. Soc.

1, 109 (1964).

[14] L. W. Barr and J. N. Mundy, Diffusion in Body-Centered

Cubic Metals, p. 171 (American Society for Metals, Metals Park, Ohio, 1965).

[15] J. R. Manning, Phys. Rev. 139, A2027 (1965).

[16] J. R. Manning, Can. J. Phys. (to be published).

[17] See for example the review article, R. J. Friauf, J. Appl.

Phys. 33 (Suppl. 1), 494 (1962).

[18] R. E. Howard and A. B. Lidiard, J. Phys. Soc. Japan 18

(Suppl II), 197 (1963); Reports on Progress in Physics

27, 161 (1964).

[19] A. B. Lidiard, Phil. Mag. 46, 815 (1955); 46, 1218 (1955).

63

Chemical Diffusion Coefficients for Some Nonstoichiometric Metal Oxides

J. Bruce Wagner, Jr.

Department of Materials Science, Northwestern University, Evanston, Illinois 60201

The chemical diffusion coefficient, 5, is the proportionality constant in Pick's law when the dif- fusion process occurs under a concentration (activity) gradient.

The chemical diffusion coefficients may be obtained from: (1) data of parabolic oxidation of a metal when certain stringent boundary conditions are met, (2) measurement of the change in weight

as a function of time for an initially homogeneous oxide crystal which is exposed isothermally to an

oxidizing or reducing atmosphere, (3) measurement of the electronic conductivity as a function of time under the same conditions as in (2), and (4) measurement of a color change as a function of time under

the same conditions as in (2). Available values of D in cm^/sec using methods (2) and (3) are:

Omho^IO-" (1150 °C) Dpeo =74.1 exp [-15.36 (0/Fe)] (1100 °C) Dcoo =4.33 X 10-3 exp (-24,000)//?r (800-1100 °C)

Dtiio = 7.52 X 10-^ exp (- 21 .900)/Rr (800-1 100 °C)

Under the assumption that the self diffusion coefficient of the anion is much less than that of the

cation, values of the cation tracer diffusion coefficient have been calculated from the chemical diffusion

coefficient using the analyses of C. Wagner [Z. physik Chem. B32, 447 (1936)] and of Darken [Trans. AIME 175, 184 (1948)]. These calculated values of tracer diffusion coefficients are in good agreement

with experimental values reported in the Uterature.

Key Words: Chemical diffusion, electronic conductivity, self diffusion, tracer diffusion.

1. Introduction

Diffusion processes occurring under a chemical

potential gradient are important in the preparation

of semiconductors, oxidation of metals, reduction

of ores, reactions in the solid-state, and sintering

of oxides. In the present context, the chemical diffusion

coefficient is important in (1) materials characterization, i.e., it allows a calculation of the time necessary to equilibrate an oxide crystal with a given

partial pressure of oxygen and (2) in certain cases

the data can be used to help verify a particular

lattice defect model in the crystal.

Examples are presented in which a nonstoichiometric oxide crystal is heated in a given oxygen

partial pressure. The oxygen gas is removed from

the experimental apparatus containing the crystal,

and oxygen gas of a different partial pressure is admitted. So long as the crystal is stable in both

these atmospheres, oxygen will be removed or added to the crystal as the case may be and for

sufficiently thick samples, the addition or removal

of oxygen will be controlled by diffusion. The

proportionaHty constant in Pick's law for such

processes will be termed the chemical diffusion

coefficient and denoted by the symbol, D. It is sometimes termed an effective or interdiffusion

coefficient.

2. Definition of Terms

We begin by defining some terms used in the

present paper. The term, self diffusion coefficient,

Z)f, means here the coefficient in Pick's law when

diffusion of the I'th species in a crystal of which i

is a component takes place in a crystal of constant composition and no chemical potential gradient

of any of the components. Thus,

Df

lim

(dcildx)-

Ji _{dCildx)\c,

(1)

where concentration, c, is used in place of activity, Ji denotes the flux and x is the distance into the

crystal in the case of one-dimensional diffusion. The particles are assumed to move independently,

i.e., cross terms in the flux equation are neglected.

The self diffusion coefficient is in turn related to

65
the tracer diffusion coefficient, Df, through a cor- relation factor,/, by

DT = fDf (2)

where / ^ 1. Thus, tracer diffusion refers to diffusion

of a tracer in a crystal of constant composition

under no chemical potential gradient. Of course

tracer diffusion can be and has been measured in crystals having a chemical potential gradient.

However, the above terminology [eqs. (1) and (2)] seems generally accepted. The self diffusion coeffi- cient of the ith species, Df , is proportional to the

diffusion coefficient for the defect responsible for the migration of the ith species. In the present

paper only oxides with majority lattice defects as cation vacancies will be considered. Hence

Df=XviDvi (3)

where Xy^ denotes the mole fraction of vacancies

and Dvj denotes the diffusion coefficient of

vacancies.

In the present context, the process of chemical

diffusion refers to diffusion which occurs in a concentration gradient or to the flow of a "deviation

from stoichiometry" [1] ' through a crystal. The

total flux J , is given by the algebraic sum of the

two fluxes in a binary compound as

• J = Jl+j2 (4)

where the subscripts 1 and 2 denote the cation and

anion, respectively. The chemical diffusion coeffi- cient, D, is the proportionahty constant in Fick's

law,

J = bidcldx) (5)

where c is the absolute value of the concentration

of the excess of one of the two components ex- pressed in appropriate units. For example, for a metal deficit compound such as CoO, c represents

the excess of oxygen above that of the stoichiometric

compound expressed in equivalents per cubic

centimeter.

3. Methods of Determining the Chemical Diffusion Coefficient

The chemical diffusion coefficient can be deter- mined by (1) the isothermal parabolic oxidation of

a metal when stringent boundary conditions are

' Figures in brackets indicate the literature references at the end of this paper.

met; (2) the change in weight of a crystal as a

j function of time when a crystal is initially equil- , ibrated in a given atmosphere of oxygen and then

j the partial pressure of oxygen changed to a new ! value; (3) the change in electrical conductivity as j a function of time for conditions as in (2); and (4) i the movement of a color front in a crystal as a function of time for conditions as in (2). Data obtained from the parabolic oxidation kinet- j ics can be used to determine the self-diffusion i coefficients as well as chemical diffusion coeffij cients. This method requires that the oxide grow- | ing on a metal be compact, exhibit no porosity,

\ and maintain good contact with the metal, Moore I [2, 3] has used this method to obtain the chemical j diffusion coefficient in cuprous oxide and recently ] Pettit [4] has summarized the relation between \ activities and diffusion coefficients in a p-type oxide ; growing on a metal. ' The thermogravimetric [5] and electronic conj ductivity [6, 7] methods of measuring the chemical

| diffusion coefficient are based upon the measure- ment of a property of the oxide crystal as the i gradient in composition or the gradient in the i deviation from stoichiometry is changed. The \ crystal is initially brought into equilibrium with a ; given partial pressure of oxygen. A different partial ] pressure of oxygen is introduced to the crystal at | time zero. The surface of the crystal is assumed to j come immediately into equilibrium with the gas I phase, i.e., the chemical potential of oxygen at the i oxide surface and gas phase are equal. A physical i parameter is measured as a function of time as

| the crystal attains its new equilibrium state. In | the case of a weight change, the solution to Fick's

j second law for a thin plate is 1

log(l-Am/A«;) = log(^)-|^ (6) 1

where Am denotes the weight change of the crystal ! at any time, t; Aw denotes the weight change at the

j new equilibrium state (i.e., the total change in i weight at the end of the experiment); and a is the

j half thickness of the slab. The equation for a brick- } shaped sample of dimensions 2a, 26, 2c is,

j

log (1 — AmiAw) = log

In eqs (6) and (7) D is assumed to be independent

of composition. The appUcabiUty of this assumption

will be discussed below.

66
When certain boundary conditions are fulfilled,

the measurement of the weight change yields an unambiguous result for the chemical diffusion coefficient. For oxides where the change in weight is small, i.e., where there are small deviations from

stoichiometry as is the case for NiO, CoO and MnO,

it is more convenient to measure the electronic

conductivity. This method is suitable when there

is predominantly only one type of carrier, electrons

or holes, and neither the mobihty of the carrier nor the degree of ionization of lattice defects

changes over the composition range of interest. In such cases, the ratio, A/n/Aw in eqs (6) and (7) may be replaced by the ratio of the corresponding

electronic conductivities Acr/AcTac, as will be shown

below.

Equations (6) and (7) are the solutions for diffusion

into or out of a finite crystal. A useful cross check

is to treat the initial kinetic data. Am or Act versus time for the beginning of the reequihbration, as

diffusion into a semi-infinite crystal Consider the

solution for a slab. The initial kinetics before the center of the crystal changes composition, can be

represented (5) by.

kpt

where kp is a parabolic rate constant and

4/)

TT

(8)

(9)

In eq (9), Ac is the total change in cation concentration for a given experiment.

For the thermogravimetric method, an automatic

recording balance is employed. For the electronic

conductivity method, the crystal is wrapped with

four platinum wires, two for current leads and two

for the IR drop. A conventional d-c measuring cir- cuit is used. The gases may be mixtures of CO and

CO2 or argon and oxygen metered in known amounts

by means of capillary flow meters. In order to minimize thermal segregation, a minimum linear flow

velocity of 0.9 cm/sec is recommended [8]. Initially the crystal is equihbrated at a given temperature and oxygen partial pressure. The gas flow

is terminated and a new gas mixture of a different oxygen partial pressure is admitted. The time when

the new gas front reaches the crystal is taken as time zero for the experiment, i.e., when the sample

just begins to change weight or conductivity.

Weight changes or conductivity changes are taken as a function of time until the new equilibrium state is attained.

The movement of a color front as a function of

time in single crystals of nickelous oxide has been

used by Slack [9] to obtain the diffusion coefficient of cation vacancies. A detailed description of this experiment is not available.

It is the purpose of the present paper to discuss

available data for a series of nonstoichiometric

oxides, MnO, FeO, CoO, and NiO. Emphasis will be placed on methods (2) and (3) listed above.

4. Relation Between the Chemical Diffusion Coefficient, Self-Diffusion Coefficient and Tracer Diffusion

Coefficient

The relation between the chemical diffusion coefficient and the self-diffusion coefficients was first formulated by C. Wagner [10], especially for the

understanding of oxidation of metals. However, his theoretical analysis was not restricted to the oxidation of metals. His equation for a binary compound

senjiconductor that is predominantly an electronic

conductor and that exhibits only small deviations

from stoichiometry is.

Z) = {|2i|Df+|z2|£>f}

c

1 dix.Me RT dine (10)

where the subscripts 1 and 2 denote respectively

the cation and anion, z denotes the valence of the

indicated species, Cgq denotes the total concentration of cations or anions in equivalents per

cubic centimeter (the molar volume times the number of equivalents per mole), c denotes the absolute

value of the excess concentration of the metal or metalloid in the same units, fxme denotes the chemical potential of the metal in ergs per gram atom

and the other terms have their usual significance.

Darken [11] derived a similar expression particularly for binary metallic alloys. It is.

D= [X2Df+XiDI] 1 +

dlnXi (11)

where the symbols 1 and 2 denote the respective

components of the binary metallic alloy, y denotes an activity coefficient and X denotes the concentration of the indicated species in terms of mole

fraction. More recently, Brebrick [1] has derived a more

detailed equation applicable to a binary compound

semiconductor that exhibits Schottky disorder and

is in the "exhaustion region." His equation is.

r /a2+4a:ai'2

(12)

67
I where the subscripts 1 and 2 denote the cation and

anion, respectively, T denotes a jump frequency, X

denotes a jump distance, S is the fraction of va- cancies that are cation vacancies, A denotes a deviation from stoichiometry and is equal to the

difference between the cation and anion concentrations, Ks denotes the Schottky constant and n, the

intrinsic carrier concentration. Brebrick's equation

has been successfully applied to lead sulfide. Unfortunately, this equation has not been tested for oxides presumably due to the lack of necessary

parameters such as the value of the lattice disorder constant or the intrinsic carrier concentration.

Moore [2] has discussed this equation with application to cuprous oxide but the lack of the previously

mentioned parameters allowed only an order of

magnitude calculation.

4.1. Limiting Case for a Predominantly Electronic Conductor Which Has Cation Vacancies as Majority Lattice Defects and

< D-L When i)§ Dfff,, one term in eq (10) or (11) may

be neglected. This approximation appears to be

valid for the oxides considered here. The equation

of Darken becomes.

[X.,DU [l + dlny,ldlnX,]. (13)

It should be emphasized that this equation was de- rived for metalhc alloys and not for ionic or partially

polar semiconductors. The equation due to C. Wagner which was derived for semiconductors becomes.

D

c

1 dfx„

RT dine

(14)

where

j

i

A^ = Avogadro's number, i e= elementary electronic charge in esu,

'

/Li,ue= chemical potential of the metal at the inner

j (superscript i) and outer (superscript a) \ surfaces of the oxide, respectively, '

^ = transference number, 1 (T = total electrical conductivity,

| z— valence, j

j and the subscripts 1, 2, or 3 denote the cation, i anion or electron, respectively.

| According to C. Wagner [10], eq (15) may be

' written in differential form as, i

300 (t^ + t-zjtsa 1 96,500 Ne

dlJi\ dcdc

dx

(17) I

where c is the absolute value of the excess concentration of metal or metalloid expressed in equivalents per cm^. Thus the expression in the curly

brackets in eq (17) is the phenomenological chemical

diffusion coefficient, D, in Pick's first law. In the case of the oxides considered in this paper,

fa ~ 1. Furthermore the chemical potential of the

metal is related to the chemical potentials of the

cations and electrons by.

(18)

Under the assumption that the particles migrate

independently, the partial conductivity for each

species, f,cr = cri, can be related to the self-diffusion coefficient by the Nernst-Einstein equation.

Equation (10) and therefore eq (14) were derived on

the basis of the tarnish layer theory for the rate of

parabohc oxidation of metals. In this case the

amount of oxide in equivalents, n, formed per unit

time, t, per unit area, A, was

(15)

where J is the flux. Ay is the oxide layer thickness

and kr is the rational rate constant in equivalents

per second for a layer one centimeter thick. C. Wagner showed that:

kr =

300

96,500 Ne

iti + t2)t3 a- (16)

^ The Kroger-Vink notation is used. The subscripts refer to the indicated lattice

site, super primes denote the effective negative charges, super dots denote effective positive charges relative to the normally occupied lattice site and ® and © denote

electron holes and electrons, respectively.

Df = BiRTIN (19)

where fi, is the mobihty and given by C. Wagner

[10] as.

Br

300

96,500 \zi\eci (20)

In this way the expression in curly brackets,

D, of eq (17) can be related to self-diffusion

coefficient. To a first approximation, the majority lattice defects in the oxides considered here are cation va- cancies. The equation^ is.

V2 o.>(,/)=Oo+r^-+,

Hence eq (3) becomes,

(21)

(22)

68
with the state of ionization not denoted. The dif- fusion coefficient of the vacancies is.

(23)

and the mobility of the cation vacancies, i?i is given [10] by eq (20). For this limiting case where

in addition to the assumption that Df, < D'l,g, the

activity of the vacancies can be set equal to their concentration (dilute solution), eq (10) becomes

D=(l+z,)jDT 1 Me VAv

(24) Me

where the concentration is now expressed in terms

of mole fraction ^v^^- The term (1 + zi) is due to the diffusion potential, i.e., the ions are accelerated

by the more rapidly moving electrons (as electron

holes) and conversely, the electrons are decelerated.

5. Experimental Results and

Discussion

5.1. Wustite, FexO

Wustite is particularly well suited for the deter- mination of the chemical diffusion coefficient be- cause the phase field is wide, and the thermodynamics of the system are well known [8]. See figure 1. The thermogravimetric technique was used by

Levin and J. B. Wagner [5]. Single crystals were

not available. Rectangular sheets of iron, about

0.3 by 0.5 by 0.02 to 0.05 in where oxidized completely to wustite. The resulting grains were colum-

\\® \® I© 1®"

-.2 0 .2 Log (Pcoj/Pco)

Figure 1. Phase diagram for the wustite phase field. The dashed lines are lines of constant composition. In terms of the atomic ratio of oxygen to iron, 0/Fe, these are: A. 0/Fe= 1.050

B. 0/Fe= 1.060

C. 0/Fe= 1.075

D. 0/Fe= 1.100

E. 0/Fe= 1.125 The dashed-dot line represents the composition of wustite when heated at equilibrium

with a constant oxygen pressure (6.83 X lO"'"* atm).

nar and the lateral grain size diameter was about

0.3 cm. The wustite was oxidized or reduced within

the Umits of the phase field. Typical plots of raw

data for reduction of undoped wustite from an oxy- gen to iron atomic ratio (0/Fe) of 1.125 to 1.050 at several temperatures are shown in figure 2. These

data are replotted in accord with eq (8) in figure 3. From the initial slopes a value of D was obtained.

These same data are replotted in accord with eq (6) in figure 4. Again D was evaluated. The values of

the chemical diffusion coefficients obtained from

eq (8) were labeled D, and those from eq (6) labeled

Di. Values for several different reduction increments are shown in table 1. The agreement between

the two sets of values, A and A, is good. In this series of experiments, the initial composition of the

wustite was always 0/Fe= 1.125. The final composition was 0/Fe= 1.100, or 1-075 or 1.050. For a given

temperature the values of D decrease as the 0/Fe

ratio increases, i.e. Z) is a function of composition.

In order to establish this compositional dependency

Hembree [12] carried out a series of experiments at

1^

2 4

6

I

°

cP o o o o o <;)

AA

A

AA

oo

A O

o

Aq o

•

oA •

• 900°C

O 950

1000

A 1050

6 Time

16 10" (sec

Figure 2. Typical plot of data for the reduction of wustite from

an initial OjFe ratio of 1.125 to a final ratio of 1.050. After Levin and J. B. Wagner [5].

69
O

O

O •

a •

o •

• • 900°C

n

• O 950

• 1000

2

•

A 1050

1100

15 30

Time • 10'^ (sec )

45

Table 1. Values of the "Parabolic Rate Constant," Kp, and the chemical diffusion coefficients from the initial, Dj, and later D, stages of the reduction of undoped wustite for various experi- mental conditions [5]

Temper- K, Di D,

ature, °C g- per cm''/sec cm^ per sec cm^ per sec

(A) Reduction of undoped wustite, 0/Fe = 1.125 to 1.050 Ac = 0.0885 g per cm^

1100 3.11 xlO-« 3.1 X lO-** 3.3 X 10-" 1050 2.43 X 10-« 2.4 X 10 '' 2.6X 10-« 1000 1.27xlO-« 1.3X lO-ii 1.4 X 10-« 950 8.21 X 10-» 8.2 X 10-' 7.0 X 10-' 900 3.85 X 10-3 3.8x10-7 3.5 X 10-'

(B) Reduction of undoped wustite, 0/Fe = 1.125 to 1.075 Ac = 0.0624 g per cm^

1100 9.22xl0-« 1.9X 10-« 2.5 X 10-« 1050 6.97 X 10-« 1.4X 10-s 1.3 X 10-6 1000 4.19X 10-'' 8.4 X 10-7 8.0 X 10-'

9.50 2.77 X 10-3 5.6 X lO-'' 5.6 X 10-' 900 1.55x 10-3 3.1 X 10-' 3.1 X 10-'

(C) Reduction of undoped wustite, 0/Fe = 1.125 to 1.100 Ac = 0.0296 g per cm^

1100 9.81 X 10-'» 8.8 X 10-' 8.6 X 10-' 1050 6.13x 10-'" 5.5 X 10-' 5.7 X 10-' 1000 3.85 X lO-'o 3.5 X 10-' 3.8 X 10-' 950 2.50X 10-'" 2.2 X 10-' 2.7 X 10-' 900 1.37 x 10-'" 1.2 X 10-' 1.8 X 10-'

Figure 3. Parabolic rate curves for reduction of wustite from OIFe= 1.125 to 1.050. After Levin and J. B. Wagner [5].

3 6 9 Time 10" '(sec )

Figure 4. Plot of data for reduction ofwustitefrom OIFe= 1.125

to 1.050 according to eq (6). After Levin and J. B. Wagner [5\.

1100 °C whereby the reduction increments were

0.01 in terms of the 0/Fe ratio. The phase field of

wustite was divided into ten increments from 0/Fe = 1.150 to 0/Fe= 1.050. See figure 5. Between a composition change of 0.01 in 0/Fe ratio, the chemical diffusion coefficient was assumed to be constant. Figure 6 shows a plot of log D versus the 0/Fe ratio obtained by P. Hembree [12]. The values of D at 1100 °C obey an equation.

/)=74.1 X

10-6-67(O/Fe)= 74 Jp-15.36(0/Fe)_ (25)

Over the entire phase field, the value of D varies

by a factor of four. Thus an oxidation experiment

and a reduction experiment, both within the phase

field from the same initial composition but to differ- ent final compositions, yield different results. Most

of the previous work involved composition ranges

which spanned one-third to two-thirds of the entire

phase field, a difference in 0/Fe ratio of 0.03 to 0.06. During the process of equihbrating a sample,

the crystal spends about two-thirds of the time of the experiment at a composition near the final composition [12]. Thus, two experiments on two crystals having identical initial compositions in the center of the phase field, one being oxidized to a higher 0/Fe ratio and the other being reduced to a lower 0/Fe ratio will give different values of b. The reverse situation whereby the same two

70
1300

1200

- 1100

iiJ 1000 2

900

800

700

050 1.060 1.070 1.080 1.090 l.lOO I.IIO 1.120 1.130 1.140 1.150 1.160

-2 0 *Z +4 LOG (Pco,/Pco) X 10*'

Figure 5. Wustite phase field with lines of constant composition (dashed lines). The numbers at the top of the dashed lines are the atomic ratios of oxygen to iron. After Hembree [12]

9.0

8.0

70

S ^°

E" 5.0 +o

X 4.0

1.0

T 1 1 r constructed a "hard sphere" model for cation

vacancies in wustite which explains qualitatively

the decreasing D with increasing vacancy concentration at high vacancy concentrations.

Landler and Komarek [14] explain a similar

observation for the reduction of wustite in H2-H2O

mixtures by a change in the mobiHty of the vacancies

as a function of composition.

ilOO°C 1050 1000 950 900

-5.5

I.IIO 0/Fe Figure 6. Compositional dependence of D for undoped wustite as a function OlFe ratio at 1100 °C according to Hembree [12].

crystals now are brought back to the same initial composition wiE yield approximately the same

value of D.

In a dilute solution of cation vacancies in a lattice,

the cation vacancy can jump to any filled cation

lattice site. When the concentration of cation

vacancies is four to eleven percent of the total cation sites, then each vacancy has some neighbors

as vacancies. The vacancies have a charge relative

to the lattice and are repelled from one another.

For such high concentrations of vacancies, the

migration of vacancies is impeded by one another.

Thus it is understandable that the chemical diffusion

coefficient which is proportional to the vacancy

diffusion coefficient should decrease with increasing vacancy concentration. P. Hembree [12] has

-6.0

-6.5

-7.0

1 1 I

O/Fe Rotios • 1.125 to 1.050

1.125 to 1.075

1.125 to l.lOO

1

1 .72 .76 .80 Temperature, lO'/T (°K"'

.84

Figure 7. Temperature dependence of the chemical diffusion

coefficient of wustite for several composition increments ac- cording to Levin and J. B. Wagner [5].

71
The temperature dependence of D for Fej-0 has

thus far only been evaluated for large reduction

increments in CO-CO2, viz. 0/Fe from 1.125 to 1.050, 1.075 or 1.100 as the case may be. An Arrhenius plot for Di is shown in figure 7 and values of

the corresponding "activation energies" are composite quantities containing both the energy for motion of the cation vacancies and the partial molar

enthalpy of dissolution of oxygen from CO..

Landler and Komarek [14] report a value of 38

kcal for the activation energy for reduction of un- doped wustite in H2-H2O mixtures. Their value was not a function of the wustite composition. They

considered the rate controlling steps to involve

the combined processes of solid state diffusion

and a phase boundary reaction at the gas-oxide

surface.

5.2. Relation Between D and D'^^.

The value of D for Fcj-O is easily obtained and

relatively unambiguous. The equation (24) of C. Wagner relating D and DJ was made expressly for compounds exhibiting small deviations from

stoichiometry and Darken's equation (13) was not derived for partially polar compounds. Nevertheless,

both equations provide an order of magnitude

agreement in relating the experimentally deter- mined values of D with experimentally determined

values of Z)^^. In table 2 are values of D for wustite

calculated using the tracer diffusion data of Himmel,

Mehl and Birchenall [13J and the thermodynamic

data of Darken [8]. The metal-metal oxide boundary was taken as a standard state for the oxygen

activity. Both the limiting equation of Wagner

[eq (24)] and the equation of Darken [eq (13)] were 1 tested. These values are to be compared with ' experimental values obtained by Levin and J. B. : Wagner shown in table 3. The agreement is good,

j expecially so when one realizes that wustite exhibits I super-lattice lines by x-ray studies which suggests I that the vacancies are ordered into "clusters" [15], | and di p \o n transition is found in the center of j the phase field [16, 17]. | Using the more recent data of Hembree [12] for ' D at 1100 °C, and taking 0/Fe= 1.07 as an average

value, I

\ Z)(expt'l)= 74.1 X 10-6-67(O/Fe)

jJ = 5.3 X 10-6 cmVsec. (26) \

As a comparison one may use equation (24) of i C. Wagner and the tracer diffusion data of Hembree

[18] and calculate the chemical diffusion coefficient i using eq (25). According to Hembree, ;

^Fe 1.1 X 10-^cm-/sec (27)

at O/Fe = 1.07 and 1100 °C. Because the iron sub- ; lattice is f.c.c, as a first approximation /=0.78, j and 2i = 2. Thus,

Z)(calc'd)=3

f X, = 5.5 X 10-6 cmVsec (28)

where A'^ is Avogadro's number, / is the correlation

' factor=0.78, Vm is the molar volume = 11.84 | Table 2. Calculated values of the chemical diffusion coefficient for wustite [5]

0/Fe No Temp. °C

Himmel et al. [13] From Wagner, eq(24) From Darken, eq(13)

D^.^, cm^ per sec D, cm^ per sec

L dlnNfJ

D, cm^ per sec

1.050 0.512 1100 =9.1 X 10-« 7.2X 10-« 58.2 3.5 X 10-6 1000 3.6 X 10-8 2.8X 10-« 57.9 1.4 X 10-6 900 2.0 X 10-8 1.5 X 10-6 51.0 6.7 X 10-6

1.075 0.518 1100 <^2.0x 10-' 10.3 X 10-" 58.2 7.7 X 10-6 1000 4.9 X 10-8 2.6 X 10-« 57.9 1.9 X 10-6 900 2.6 x 10-8 1.3 X 10-6 51.0 0.9 X 10-6

1.100 0.524 1100 "=2.2 X 10-' 8.6 X 10-6 58.2 8.6 X 10-6 1000 7.4 X 10-8 2.8 X 10-6 57.9 3.0 X 10-6 900 3.6 X 10-8 1.4 X 10-6 51.0 1.2 X 10-6

* These values are to be divided by 0.78 to yield the self-diffusion coefficient. By extrapolation.

72
Table 3. Experimental results obtained from the Arrhenius

curves for the reduction of undoped wustite [5]

Activation Activation Experiment energy from energy from

ln(T lug U]n ve 1LIIT1 y log Lfl VS. 1/ J , kcal per mole kcal per mole

Reduction of undoped FeO: 0/Fe= 1.125 to 1.050 34.0 + 2.4 37.4 + 3.1 0/Fe= 1.125 to 1.075 29.2±2.9 32.1 + 1.4 0/Fe= 1.125 to 1.100 31.5±0.4 24.8 ±0.7

^ The errors represent 80 percent confidence limits.

cm^/mole, and [V'^g] is the vacancy concentration

per cm'' ~ 4 X 10^^ The agreement is remarkable

in view of the assumption of an ideal solution and

the assumption that /= 0.78. In the studies by Levin and J. B. Wagner [5], by Landler and Komarek [14] and by Hembree [12],

D decreases as the 0/Fe ratio (which is proportional

to the cation vacancy concentration) increases. Al- though a calculation such as that presented by eq

(28) yields the correct order of magnitude, if values

of the tracer diffusion coefficient are calculated

using eq (13) from the experimental values of Z), the calculated values of D^^ also decrease with

increasing 0/Fe ratio [14]. This is not in accord

with radiotracer experiments which show that values of D^^ increase with increasing 0/Fe ratio

[13, 18, 31]. Landler and Komarek [14] have dis- cussed this problem in terms of the mobihty of

cation vacancies which is not a constant but a function of the composition.

volume and this in turn is related to the change

in the electronic conductivity. Act, by

Unit volume

A[Oo] = A[( -~1 Ao-

(30)

where e denotes the electronic change and u, the

mobihty of the electron holes. Then using the subscripts a and b for an initial and final states.

iWn

(31)

If the electron hole mobility is independent of

oxygen pressure over the pressure range in question,

then the ratio Ubiua will be unity. The conductivity

is written as

o- = [©] eu = K'pH'^ exp

ag;/2-ag|-

RT

and the concentration of electron holes is

[®]^K"p}i' exp (-AG}I2RT).

(32)

(33)

The terms, AG^ and AG| are the free energies of

dissolution of oxygen in the crystal and that for motion of the electron holes respectively and K' and K" are two constants dependent only on temperature. The electron hole mobility may be ex- pressed by

u = constant • exp RT R

(34)

5.3. Cobaltous Oxide, CoO

Cobaltous oxide has a narrower range of homogeneity than does wustite. The maximum vacancy

concentration is about 1 percent at 1000 °C [19]. The chemical diffusion coefficient could be obtained

from a thermogravimetric technique, but the d-c electronic conductivity was more convenient. In order to use this technique, the ratio of the conductivity change at any time t to that at infinite time, AalAcr^, must be substituted for the ratio Am/Aw as was discussed previously. In this case

the evaluation of D requires that the mobility of the

carrier remains constant and the particular model

for the point defects in the crystal must be known.

In the case of CoO, the model has been shown to be [19, 20],

V2 02(g) = Go +F^o+®- (29)

Consequently, a weight change per unit volume is equal to the oxygen pickup (or evolved) per unit

where AG^^ A//|— TAS®. Tannhauser and Fisher

[20] report that both u and A//© are independent

of oxygen pressure for the experimental conditions

reported here. Furthermore, Eror [19], found that AS ^ also was independent of oxygen pressure. These

data suggest that u is independent of the oxygen

partial pressure and therefore AcrjAcTy, may be

substituted for AmIAw in eqs (6) or (7). The fore- going model depends on the hopping process for the movement of the electron holes. If this model

is not valid, the substitution may still be used

provided that the mobility ratio is unity.

Equation (7) as modified was used by Price and

J. B. Wagner [7] to determine D for CoO at 800,

900, 1000 and 1100 °C. The data may be expressed

by

D = 4.33 X 10-3 exp (- 24,000//?r) cmVsec. (35) The relation between D and Z)^^ was obtained using

the thermodynamic data of Eror [19] and the equa290-621 0-68—

6

73
Figure 8. Cation vacancy concentration in cobaltous oxide as a function of oxygen pressure and temperature according to Eror [19].

tions of Darken or C. Wagner. The thermodynamic

data of Eror are shown in figure 8. The metal,

metal-oxide phase boundary [21] was used as the

reference state for the oxygen activity. A comparison

of the calculated values of D"^^ and the experimental

data are shown in table 4. Again the agreement is satisfactory.

5.4. Nickelous Oxide, NiO

Of the series of oxides under consideration, nickel

oxide exhibits the smallest deviation from stoiTable 4. Comparison of calculated and experimental values of Dfo (cm^jsec) at one atmosphere of oxygen [7]

Temper- Price [7] Calculated Experimental

ature °C eq (13) eq (24) Carter and

Richardson [24]

1100 13.9 X lO-'* 1.54 X 10-" 6.50 X 10-" 1000 4.07 X 10-" 0.62 X 10-" 6.50 X 10-" 900 0.906 X 10-« 0.22 X 10-" 0.742 X 10-" 800 0.391 X 10-" 0.06 X 10-" 0.180x10-9

chiometry, about 0.1 percent at 1000 °C. It was assumed that because NiO has the same structure

as CoO that the conductivity method could be applied. Price and J. B. Wagner [7] determined

b for NiO between 800 and 1100 °C to be

D = 7.52x10-4 exp (-21,900//?r) cm^/sec. (36)

For the calculation of the tracer diffusion coefficient, the data of Mitoff for the concentration

of cation vacancies in NiO was used. These data were extrapolated to the lower temperatures of

interest in the present work. At 1300 °C Mitofif [30]

obtained

2atm_

0.2 atm

1/6 6X lO-'^ exp (-ll,400)/«r

vac/ion pair. (37)

The data of Eror [19] suggested that singly ionized

vacancies may be the majority defects in NiO be- tween 800 to 1100 °C. Therefore the equation of

Mitoff may be modified to yield

J^atm_

0.2 atm

1/4 6 • 10-3 exp (-11,400/fir)

vac/ion pair. (38)

The mole fractions of oxygen and nickel were cal- culated from this equation. The standard state for the activity of oxygen was the metal oxide phase

boundary as determined by Kiukkola and C. Wagner [21]. Values of D^^ were computed using the

Darken equation. The use of the limiting equation

(24) of C. Wagner, gave lower values of Z)^j. On

using the more general equation (14) of C. Wagner

with no assumption of ideahty, the values of D''^^ were very nearly the experimental values of tracer

diffusion. See table 5.

Table 5. Comparison of calculated and experimental values of DJj [cm'^jsec) at 0.209 atmospheres of oxygen [7]

Temperature °C 1100 1000 900 800

Price and Wagner [7] eq (13) Price and Wagner [7] eq (25) Linder and Akerstrom [25] Fueki and Wagner [26] Shim and Moore [27] Choi and Moore [28]

3..59X 10-"

0.80 X 10-"

3.31 X 10-"

0.842 X 10-"

3.60 X 10-"

9.55 X 10-"

1.12 X 10-"

0.22 X 10-"

0..55X 10-"

0.198 X 10-"

0.999 X 10-"

2.49 X 10-"

0.429 X 10-"

0.08 X 10-"

0.036 X 10-"

0.036 X 10-"

" 0.227 X 10-"

0.534 X 10-"

0.077 X 10-"

0.02 X 10-"

0.005 X 10-"

0.005 X 10-"

" 0.039 X 10-"

" 0.085 X 10-"

By extrapolation.

74
j 5.5. Manganous Oxide, MnO

Manganous oxide exhibits an apparent p to n

transition in its phase field [19]. See figure 9. Therefore conductivity measurements can only be

used for an evaluation of D over those composition ranges where the mobilities of the predominant

I carriers are independent of oxygen pressure and

I only one type of carrier is responsible for the

majority of the electronic conductivity. It has been

! assumed that the mobilities of the electrons or of

the holes are constant at oxygen partial pressures

,j near either extreme of the composition range, e.g.,

l| at 1150 °C for p-type, 10"" ^ Pq, (atm) ^ IQ-^, and

chromium, or with hthium has been studied. As a

first approximation it is assumed that the majority

lattice defects in wustite are doubly ionized cation

vacancies. Accordingly the following models hold

TiO2 = Tip^ + 20o+r;, (39)

Cr2a = 2 Crp^+30o + r;^ (40)

Li.20 = 2 Lij;^+0o + 2e. (41)

On an atom for atom basis, titanium should cre- ate twice as many vacancies as chromium, each for

E

o

e

oD

10' -

id"2 I I

'

I I I I I I I I I I L_

10^" ,-20 ,^19 10'^ lO'a ,^ia ,^17 ,^I6 ,0'° ,^I5 10'-' 10'^ ,^13 10'-^ 10'^ 10" ,^10 10'" ,^9 .nS ,^7 10 =

Pq^ (atm.)

Figure 9. Electronic conductivity of manganous oxide as a function of oxygen pressure at temperature increments of 50 °C- The lowest curve corresponds to 900 °C. After Eror [19].

for R-type, 10""* ^ Pq^ (atm) ^ 10-^\ Thus far tentative values for p-type MnO yield D ~ 10~^ cm^/sec

at 1150 °C [22]. The temperature dependence of

this value over the temperature interval of 850 to 1150 °C is very small. The reason for this is not known, but these values are prehminary and must

be regarded as such. As a comparison, the value of has been calculated by Fueki and J. B. Wagner

[23] to be about 10"" cm'/sec for p-type MnO for 1150 °C.

6. Effect of Impurities, Intentionally

Added or Not

Wustite is the only oxide in which a systematic

study of the effects of trace impurities on D has

been made. Wustite doped with titanium, with

a given oxygen pressure and temperature. Conversely, lithium should destroy vacancies. Because

D is inversely dependent upon the vacancy concentration, the addition of higher valence solutes is expected to lower the value of D and conversely

for lower valence solutes. Levin [5] and Hembree

[12] have shown that the addition of chromium

decreased the value of D relative to undoped

wustite. Likewise titanium decreases and hthium

increases D relative to undoped wustite [12]. The

results are in agreement with the simple lattice models given by (39), (40) and (41). The samples

doped with lithium were prepared by evaporation

of metallic hthium onto a wustite slab and subsequently annealing in a CO-CO2 gas mixture. The

samples containing titanium or chromium were made by oxidizing an alloy to a given oxygen to

75
metal ratio. This procedure is subject to experimental problems due to the phenomenon of internal

oxidation during the preparation of the doped wustite. Even if there was a precipitate of another oxide

formed initially, the samples were annealed for periods of weeks and the results were clearly in accord with theory.

The nickel oxide crystals were impure. Samples

were sent to different companies for spectrographic

analyses and to one company for neutron activation analysis. The results were not consistent. Accurate and consistent values of analyses are needed

to test the role of trace impurities. It must be emphasized that trace impurities could act as traps

for electron holes and in so doing change the oxygen

pressure dependence of conductivity. Experiments

are needed to test the dependence of D, D^. and cr on added impurities in CoO and MnO as well

as NiO.

7. Remarks on Other Experimental

Errors and Precautions

The obvious precautions for measurement of

weight, of sample geometry or of probe spacing are

relatively easily made. Other precautions, in addition to an awareness of the crystal purity cited pre- viously, include a test of the assumption that the

oxide surface comes into equilibrium with the gas phase virtually instantaneously when a new gas mixture is admitted to the surface. Especially for the case of wustite in CO-CO2 mixtures, there have

been suggestions [14] that a surface reaction is rate-controlling even though the data can be fitted to diffusion controlled kinetics. The data for wustite

are self-consistent and the doping experiments

should be carried out whereby the chemical potential of oxygen in the gas phase is maintained

constant but the CO-CO2 mixture is diluted with an inert gas. In such a case, the kinetics should

remain unchanged if the reaction is diffusion controlled because the rate depends only on the gra- dient of chemical potential in the oxide. Such

studies are underway for the case of wustite. In the case of using the electronic conductivity

method at low oxygen pressures provided by oxygenargon mixtures, e.g., NiO at ^02^10"*^ atm, and

1100 °C, the voltage drop across the probes occasionally did not stabihze over a long time period

but fluctuated, e.g. by a factor of two and one half. In these cases, it is suspected that the surface is not completely equihbrated with the oxygen pressure. This phenomenon cannot be avoided so long

as this method is used. A check on the precision

and accuracy of the data can be made by comparing the final (equihbrium) value of the conductivity

to equilibrium values on plot of log (x versus poa

where the data have been extended to higher pres- sures where this fluctuation phenomenon is not observed. Likewise gas mixtures, e.g., CO/CO2,

where the ratio of one gas to the other is less than

about one to one hundred will have a very small

"buffer capacity" against small fluctuations in composition [4]

.

In the solution of Kick's second law, D was as- sumed to be constant. This assumption is better

the smaller the composition ranges over which the

data are computed as was determined for wustite.

Thus the composition dependence of D can be

determined by the evaluation of D for composition

increment which are sufficiently small so that the

error in assuming D a constant is less than the

experimental error. The relations between D and DT were derived

by C. Wagner for compounds exhibiting small deviations from stoichiometry. In such cases the

values of the valence of the cation is relatively well- defined. In case of large deviations from stoichiometry, the assumption of integer values may not be

correct. Meussner and Birchenall [29] and Carter

and Richardson [24] considered this problem for iron sulfide, and for cobaltous oxide, respectively.

For the most part other investigators have neglected

this consideration. On the other hand, the thermodynamics are best known for oxides exhibiting large

deviations from stoichiometry owing to the relative ease of the experimental procedure. These competing factors must be weighed when the equations

relating D and D] are used.

8. Summary

It has been shown that the chemical or inter- diffusion coefficient for oxygen in crystals of

manganous oxide, ferrous oxide, cobaltous oxide

and nickelous oxide can be determined by a thermogravimetric or electronic conductivity tech- nique. Values so obtained can be used to calculate

tracer diffusion coefficients in these oxides. The

calculated values are in good agreement with

experimental values reported in the literature.

These studies have been supported by the Advanced Research Projects Agency through a grant

from the Materials Research Center of Northwestern University. The author expresses his sincere thanks to helpful discussions with P. Hembree, F. S. Pettit, and J B Price.

10. References

[1] R. F. Brebrick, J. Appl. Phys. 30 , 811 (1959).

[2] W. J. Moore, J. Chem. Educ. 38, 232 (1961).

76
[3] W. J. Moore, J. Electrochem. Soc. 100, 302 (1953).

[4] F. S. Pettit, J. Electrochem. Soc. 113, 1249 (1966).

[5] R. L. Levin and J. B. Wagner, Jr., Trans. Am. Inst. Min.

Met. Engrs. 233, 159 (1965).

[6] H. Dunwald and C. Wagner, Z. physik Chem. B24, 53 (1934).

[7] J B Price and J. B. Wagner, Jr., Z. physik Chem. (N.F.)49,

257 (1966).

[8] L. S. Darken and R. W. Gurry, J. Am. Chem. Soc. 67,

1398 (1945).

[9] G. Slack, private communication to S. P. Mitoff, J. Chem.

Phys. 35 , 882 (1961).

[10] C. Wagner, Z. physik Chem. Bll, 139 (1930); Z. physik.

Chem. B32 , 447 (1936).

[11] L. S. Darken, Trans. Am. Inst. Min. Met. Engrs. 175, 184

(1948).

[12] P. L. Hembree, Masters Thesis, Northwestern University.

Paper submitted to Trans. Am. Inst. Min. Met. Engrs. May 1967.

[13] L. Himmel, R. F. Mehl and C. E. Birchenall, Trans. Am.

Inst. Min. Met. Engrs. 197, 827 (1953).

[14] P. F. J. Landler and K. L. K. Komarek, Trans. Am. Inst. Min. Met. Engrs. 236, 138 (1966).

[15] F. Koch, Ph. D. Thesis, Northwestern University (1967).

[16] I. Bransky and D. S. Tannhauser, Trans. Am. Inst. Min.

Met. Engrs. 239, 75 (1967).

[17] W. J. HiHegas and 'J.'B.""Wagner, Jr., Phys. Ltrs. 25 A,

742 (1967).

[18] P. L. Hembree, Ph. D. Thesis, Northwestern University

(1967).

[19] N. Eror, Ph. D. Thesis, Northwestern University (1965).

[20] D. S. Tannhauser and B. Fisher, J. Chem. Phys. 44, 1663

(1966).

[21] K. Kiukkola and C. Wagner, J. Electrochem. Soc. 104,

379 (1957).

[22] J B Price, Ph. D. Thesis Research in Progress, Northwestern

University.

[23] K. Fueki and J. B. Wagner, Jr., J. Electrochem. Soc. 112,

970 (1965).

[24] R. E. Carter and F. D. Richardson, J. Metals 200, 1244

(1954). [25) R. Lindner and A. Akerstrom, Disc. Faraday Soc. 23,

133 (1957).

[26] K. Fueki and J. B. Wagner, Jr., J. Electrochem. Soc. 112,

384 (1965).

[27] M. T. Shim and W. J. Moore, J. Phys. Chem. 26, 802 (1957).

[28] J. S. Choi and W. J. Moore, J. Phys. Chem. 66, 1308 (1962).

[29] R. Muessner and C. E. Birchenall, Corrosion 13, 677 (1957).

[30] S. P. Mitoflf, J. Chem. Phys. 35 , 882 (1961).

[31] P. Desmarescaux and P. Lacombe, Mem. Sci. Rev. Metallurg. LX, No. 12, 899 (1963).

77
l|' 1iIi1 ii ; i 1 J I )
The Use and Limitations of the Concept of an Effective Binary Diffusion

Coefficient for Multi-Component Diffusion

Alfred R. Cooper

Case Western Reserve University Cleveland, Ohio 44106

An "effective binary diffusion coefficient" is frequently used for treating kinetics of mass trans- port controlled reactions in multicomponent systems. The concept is examined and the relation of an EBCD to the matrix of phenomological coefficients is determined. It is shown that the EBCD are useful parameters in that Pick's laws are meaningful if:

(i) the concentration gradients of all species are in the same direction.

(ii) either a steady state exists or there is no characteristic distance involved in the problem

(i.e., infinite or semi-infinite diffusion). When these conditions are obeyed, use of the EBCD gives considerable convenience dealing with mass transport in particular cases. This convenience is gained at the expense of a loss in generality, because the EBCD depend on direction (in composition space) as well as composition.

Key Words: Composition dependence of diffusion, diffusion, effective binary diffusion coefficient.

Introduction

The kinetics of many important processes,

e.g., dissolution, precipitation, solid-solid reaction,

and homogenization depend on interdifiFusion, and

frequently technically important systems are multi-component. Fortunately, the phenomenology

of multi-component diffusion seems reasonably

well understood [1].' However, there have been

relatively few examples [2] of unambiguous deter- mination of the matrix of coefficients necessary

to describe the diffusion process in multi-component systems (possibly none in oxide systems)

and few [3] efforts to predict, by solution of the

simultaneous differential equations of multicomponent diffusion, the behavior of a non-uniform

multi-component system that is tending toward

uniformity. Typically, the determinations and pre- dictions have involved the often questionable

assumption of composition independent coefficients. Because of the complexity of the usual approach, we deem it advisable to critically examine a some- what simpler approach to multi-component dif- fusion, involving the concept (perhaps first considered by Hougen and Watson) [4] of an ef- fective binary diffusion coefficient for multicomponent diffusion (henceforth EBDC). Since the

EBDC are often used without concern as to their

validity, our goal will be to determine the relation- ship of an EBDC to the matrix of phenomenological

coefficients and the appropriateness of an EBDC

' Figures in brackets indicate the literature references at the end of this paper.

when used in Fick's equations. It is hoped that

this wiU permit a better understanding of hmitations of the EBDC's and lead to increased utility of the EBDC in appropriate circumstances.

Multi-Component Systems

To be consistent with common usage, a system

is termed ''multi-component'" if it has two or more

components that may vary independently, and

''binary' if only one component can vary independently. Thus, for example, a silver-gold alloy

would be considered a binary unless the defect

concentration can vary from its equilibrium value. Then and only then can it be considered multicomponent (ternary if there is only one species of

non-equiUbrium defect). Likewise, an oxide is a binary if it is stoichiometric and contains two

cations (although there are three species, two

cations and oxygen, electrical neutrality and the

fact that the sum of cation fractions is unity reduce

the number of species that are independently

variable). On the other hand, because it is not

stoichiometric, iron oxide, Fe(i-8)0, which has

received so much attention at this conference [5]

is also a binary oxide system even though it has a

single cation species. A different viewpoint of

Fe(i-8)0, leading to the same conclusion, would

be to consider Fe + + and Fe + + + as different species.

With these definitions, a binary system is one which has "composition space" that is unidimensional, i.e., a single number will describe the com79
position, In general, in a system of r components,

compositional space has (r— 1) dimensions as the

composition, X={xi,X2 Xr-i), is given by

a set of (r— 1) values of the fraction Xi of the in- dividual components. In what follows, the rth

species has always been chosen to be the dependent.

Where necessary the dependent species is specified

in the notation used in this paper by a presuffix, i.e.,

rDij orrFi.

A continuous "path" in composition space is given by a one-dimensional set of connected compositions. This is illustrated schematically for a ternary system by figure la. At a particular composition, the "direction", rfi of a continuous path

in this space is defined by a set of (r — 2) derivatives with respect to Xj, i.e.,

rFi = (/ii,/2i, 1)1. /(r-i),)

where

^ dxk

; -^"'^dxi

(We are not concerned with/A:;—>oo because this has expUcit meaning in terms of direction.) Various

directions, cFb^/ab, in composition space in a ternary system are shown in figure lb. Since it will be necessary to refer to positions,

paths, and directions in both real space and composition space in the following, an effort has been

made, even at the expense of awkward language, to make clear whether composition space or real space

is involved.

Multi-Component Diffusion

In 1945 Onsager [6] suggested extending Kick's

law to multi-component systems by allowing the

flux density,- y,, of species i to be a linear combination of the concentration - gradients, VpA-(A:=l^r)

of all species. Reducing this idea to a set of equations relating independent flux densities - (by establishing a reference frame) and independent fraction ^

r gradients Vxj (by noting that ^ Vxj = 0) gives

j = i what has come to be known as the phenomenological equations for multi-component diffusion. For an r component system the foflowing (/ — 1)

^ Concentration, density, fraction and flux density can be defined so as to use any

quantity: mols. number, mass, or volume to describe the amount of the species, but consistency is essential.

equations are obtained:

S'-^upVxj (1)3

for i from 1 to r — 1.

The validity of this linear combination is con- firmed by the experimental demonstration that at a given composition the coefficients rAj appear to be

independent of time [7] and direction [2d]. Another approach to multi-component diffusion

is to use the concept of an EBDC. For species i the EBDC, Z);* , defined as follows:

ji^-D,lp^Xi (2)3

is the proportionality constant between the flux density of species i and the product of density, p, and the gradient, Vx,, of fraction of i. It is clear that the EBDC, Z);^ , for species i, treats the system as if all of the other constituents can be combined together as a single species, m,

standing for mixture. Thus it is like a binary dif- fusion coefficient between i and the remainder

of the constituents. While the notion of an EBDC

arose in multi-component gas systems, it has been

used explicitly in condensed systems [8]. Conservation of species i and assumption of constant density, p, allows us to write the (r— 1) transient diffusion equations from eq 1 as follows

^=v;£,.AjVx, (3) dt j=i

i=l^(r-l).

Similar treatment yields the following equations:

^=VZ)*Vx,- (4)

i = l^(r-\)

from equation 2. We turn to the relation between the matrix of

phenomenological coefficients, rDij, and the

EBDC, Df,„. This is readily determined because combination of eqs 1 and 2, with the assumption that the

gradient of all species is in the same direction in real space, yields

D^n = E rDuidxjIdxi). (5) j=l

The reference frame for the flux densities, ^'j, must be well-defined and in what fol- lows it is assumed that the reference frame is identical for equations 1 and 2.

80
(The assumption is obviously fulfilled if we are concerned only with diffusion that is confined to a

single direction in real space. Henceforth, we confine ourselves to this situation and allow y to repre- sent distance in this direction in real space.)

Thus, in principle, the EBDC for species, i, D^„ . is calculable from (r— 1) values of ,Dij and

the set of (r — 2) independent components of the

direction vector rF, on composition space.

In addition to being functions of composition X,

the EBDC are also functions of direction rFi in composition space. This is not surprising because

in an r component system there are (r/2)((r— 1)

irreducible independent phenomenological coefficients and only (r— 1) independent EBDC.^ Thus

the phenomenological coefficients must be more

general.

A hnear (Fickean) treatment of diffusion imphes

that the coefficient, Z)f,„, is independent of the

gradient, Vj:,-. Thus for the EBDC to be useful

coefficients, in the sense that equations 2 and 4

have their usual meaning, the EBDC must be in- dependent of the gradient, i.e., they must be functions'^ of composition. Since, by definition, the

phenomenological coefficients, rA/, are functions

of composition, according to equation 5 the EBDC.

Z)jm , wiU be functions of composition if the direction, rFi, is a function of composition.

We know that in general the direction, rFi. cannot be a function of composition X because

typically there are an infinity of possible directions

(see fig. lb) at a particular composition. However,

for a specific experiment described by its initial and boundary conditions, the direction may be a function of composition. If it is and if the end

point compositions are constant, there is one and

only one diffusion path for each pair of end point

compositions. and X[^K

We wiU now consider the use of the EBDC in three different classes of experiments: (i) steady

state, (ii) transient without a characteristic length,

and (iii) transient with a finite characteristic length.

Steady State

In one sense this can be a somewhat trivial consideration, because, by definition, at steady state the diffusion path is a function only of the pair of end point compositions X^g"^ and J^'/'. Thus, when

these are fixed, the directions are a function of composition, i.e.. F = F{X) and the concept of an EBDC is valid for steady state diffusion.

* Actually, although the Onsager Reciprocal Relations [1] do not apply direcUy to the phenomenological coefficients, along with the thermodynamics they reduce the number of independent coefficients from (r— 1) - to ^r(r— 1).

^ We use function of composition in the sense that for a given value of the composition there is one and only one value of the function.

It only remains to inquire whether steady state

is ever attained in multi-component diffusion. This

is a question which needs an experimental answer. Although an unequivocal experiment (one in which

it is shown that the diffusion path remains constant)

has perhaps not been conducted, indications are

that a steady state is attained in multi-component

diffusion. Steady state in liquid systems can occur from a balance between the diffusive and convective flux densities. For a multi-component system with a

velocity, Vy, in the diffusion direction, the assumption of a steady state along with the use of eqs 3 and 4 gives:

and

dxi

(6)

(7)

Solution of these equations is important in analyzing

the kinetics of dissolution or precipitation with

free or forced convection.

To obtain the flux density of species i from eq 1 requires the solution of (/ — 1) simultaneous equations hke eq 6. It is clear that this would be a

difficult task and it is not known if there are cases

for which this has been achieved.

On the other hand, use of the EBDC requires

only a single equation 7 to be solved. In the for- tuitous case where the EBDC varies only slightly

with composition, analytical solutions [9J of equation

7 are available for conditions appropriate to several

important processes. Some support for the existence

of a steady state in multi-component diffusion is obtained from the fact that for given end member

compositions, the has been found to be in- dependent of V,/ [8a].

Transient Without a Characteristic

Length

So-called infinite and semi-infinite diffusion

couples permit the use of the "Boltzmann trans- formation," y* — y/t^'-, where y is distance in the

diffusion direction, because none of the "boundaries" are at some characteristic finite lengt i. L. Substituting this into eq 3 yields

W d r, dxj 1 ^ dxi

dy' dy* 2-' dy""

j=l^ (r-1).

(8)

81
Solution of (r— 1) simultaneous total differential equations of equation 8 has been achieved in ternary systems (r=3) for cases where the rDi/s

are constant [3| and suggestions as to how to proceed

when the coefficients depend on composition have

recently been advanced [3dJ. The existence of a solution (determined or not)

to equation 8 imphes that for given boundary conditions, the composition, X, is a function of y*,

i.e., X = X{y*). Thus, the diffusion path is a function

of the Boltzmann variable, y*.

If we restrict attention to a single phase region

(thermodynamically speaking), then there is a one-to-one relationship between y* and X as every

value of composition, X, corresponds to one and

only one value of y*. This permits the conclusion

that the diffusion path is a function of the pair of boundary conditions to equation 8" (i.e., the end

member compositions). With this estabhshed,

it follows that the direction is a function of composition, i.e., F = F{X), and therefore that the

EBDC are valid Fickean diffusion coefficients for transient unidirectional diffusion when there is no

characteristic length. Thus, rewriting equation 3

as follows

_d_ dx; , 1 dxj

dy* (9)

is permissible and solutions from heat transfer and

binary diffusion are appropriate. Because of the

simplicity of equation 9 relative to equation 8,

it is convenient to use equation 9 along with equation 2 when it is necessary to calculate the flux density of a particular species in unidirectional

multi-component diffusion in the absence of a

characteristic distance, L.

Transient With a Characteristic

Length

In so-called finite diffusion experiments where a characteristic length, L, prevents all the boundary

conditions and initial conditions being expressed

in terms of y*, eqs 8 and 9 cannot be used. The

existence of a characteristic length, L, also permits

the end member compositions X[^K A'J,"* to vary and thus the diffusion path will not be constant. Therefore, we cannot be certain that the directions,

rFj, will remain functions of composition during

the experiment and. in general, caution need be

exercised to assure that a change in end member

composition will not markedly influence the EBDC.

To consider this point, we rewrite eq 5

D*, = rDii + rfu rDn + r.k, rD,, +

•

and notice that if either the changes of direction

rFi are only very slight or if the cross terms, r^y,

/' 7^ i are very much less than the diagonal term

,Z),/, then we still may have sufficiently insensitive to factors other than composition so as to be useful. Particular care has to be exercised in cases such as illustrated in figure lb where a component of the direction can reach infinity.

Guy, Leroy and Lindemer [3d|, in considering

ternary alloys with the usual phenomenological

equations 1 and 3, have shown how, when the ratio of the magnitude of the diagonal term to the cross terms is great, homogenization of lamina of discreet

thickness, L. proceeds almost identically to that predicted using a single EBDC.

Uphill Diffusion of One Species

The EBDC for a particular pair of end member

compositions can be obtained by making an infinite diffusion experiment and using eq 9 to interpret

the results, as in terms oi D;m{X). At times," a species in a multi-component system

in a stable single phase region will behave as if its EBDC is < 0. This is sometimes termed "uphill

diffusion" [3cJ. It may occur, for example, when

the flux of a species is caused primarily by the ran- dom "hopping" of other species. It results in a

diffusion path in composition space which is S- shaped, as shown schematically in figure la, and

has regions where some of the components of the

direction,/,/}, (see fig. la) go to infinity.

For an infinite dittusion experiment we have

shown that composition is a function of y*. To be

consistent with figure 1, a graph of one component.

xk, of the composition would appear as is shown

schematically in figure 2. This profile of the xk

component versus composition requires drastic

behavior of the EBDC for species B as can be

determined by expanding eq 9 to obtain

d'-xii dDnm

«"V/y*- dy*

dxn . 1 y*dxB ^ „ dy* 2 dy* (10)

Since XH(y*) is known, it is convenient to write

equation 6 as a linear first order differential equation in Djfii,

dy*

+ piy*)D*, = ^iy*) (11)

''This conclusifin has been verified because it has been found both in multi-coni- ponenl melal |7] and oxide glass systems [10] thai the diffusion path in "infinite"

difTiision (ouples does not vary with time. [2d).A particularly clear example is found in Path 5 of figure 8 and figure 9 of reference

82
Figure 1. Ternary composition space (a) A curved diffusion path, {b) Various directions. Fab. on composition space.

whe

/^(.V*) =

d-XB /(lx,i\ 1 / 1

While eq 11 is. in principle, solvable by use of an

integrating factor, it is difficult for cases where

the concentration profile is similar to that shown

in figure 2. because as indicated by figure 3. p

(obtained from the concentration profile of figure 2) has discontinuities. An estimate [10| of D^,,, as a function of j* can be obtained which is qualitay= yt- '2 ARBITRARY UNITS

Figure 2. A schematic diffusion profile offraction of component

B, Xb, versus Boltzmann variable, y* = yt""^.

tively consistent with figure 3, with eq 7. and also with eq 3. This is shown as figure 4 and the "bad"

behavior (discontinuities, infinite values and negative values) is obvious.

The erratic behavior of the EBDC, D^^j. in special cases, limits its usefulness, i.e., for the

diffusion path in figure la it would be difficult to characterize the functional dependence of Z)j^,„ on

composition. However, two factors moderate this limitation. (1) In a ternary system there is usually

rif'i ARBITRARY UNITS

Figure 3. Schematic of the function p(y*) versus y 83
0

J,

y=yt 2 ARBITRARY UNITS

Figure 4. Schematic of DgM versus y*. {Note y* scale identical on figures 2, 3, 4).

only at most a single species in a given experiment

that has an EBDC that is erratic. The diffusion

behavior of the system can be described by choosing

the EBDC of the other two species. (2) Even when

the EBDC behaves erratically, it still follows eq 10,

and such important values as flux densities inte- grated across a given plane can be calculated.

Summary and Conclusions

The EBDC approach to multi-component dif- fusion (eqs 2 and 4), is in no way contradictory

to the usual phenomenological approach (eqs 1 and 3). In fact, when the gradient vector in real space of all species is in the same direction, the

EBDC can be derived from the phenomenological

coefficients and the direction F in composition

space. Conversely, the (r—lf phenomenological

coefficients can be determined from the directions

and the EBDC of (r— 1) experiments. The EBDC

are useful for describing mass transport in steady

state and transient behavior when the boundary

and initial conditions can be described in terms

of a variable, y*=yt~'^l-, that combines distance

and time. They may be useful in other cases as a close approximation of reality, but care has to be

taken to be sure that the EBDC are insensitive to factors other than composition before using them

in Fick's diffusion (equations 2 and 4). Using the EBDC we obtain the convenience of

having to deal with a single coefficient (the EBDC)

to relate the flux density of a species to its gradient

(eq 1) or to predict the transient behavior in a

particular infinite or semi-infinite diffusion couple

experiment (eq 10). It is much easier (1 experiment)

to determine the concentration dependence of an EBDC along the diffusion path in composition

space than it is to determine (r— 1 experiments

at each composition) the concentration dependence

of the matrix of phenomenological coefficients

along this path.

We do not obtain something for nothing. We have

gained convenience of the EBDC at the expense

of generality. Thus, while the EBDC may be very

useful for characterizing the kinetics of a process

in a multi-component system where mass transport

is rate controlling, on the other hand it would be

impractical to characterize the diffusion behavior

at a particular composition by the EBDC, because

of their direction dependence. In a ternary, for example, at a specific composition, it requires a matrix of four coefficients ^ to characterize the

diffusion behavior using the usual phenomenological

approach. Two EBDC's are required for the same

purpose. These, however, are not simple coeffi- cients, but functions of direction in composition

space, and hence they are much less efficient in characterizing the general behavior.

The help of Milton Lees, Arun K. Varshneya and

the support of NASA, Research Grant No. NCR

36-003-100, are gratefully acknowledged. Critical

comments by A. Lidiard provoked important improvements in the paper and editorial suggestions

by A. Franklin increased its readability.

References

[1] D. Fitts, Non-Equilibrium Thermodynamics, McGraw-Hill

Book Company, New York, 1962.

[2] (a) J. S. Kirkaldy, G. R. Mason, and W. J. Slater, Canadian

Mining & Metallurgical Bulletin, 64, 53 (1961).

(b) A. G. Guy and J. PhiUbert, Z. MetaUkunde, 56, 841

(1965).

(c) T. O. Ziebold and R. E. Ogilvie, Trans. AIME, 239,

942 (1967).

(d) J. P. Sabatier, A. Vignes, Memoires Scientifigues Rev.

Metallurg., 64, 225 (1967).

[3] (a) L. J. Fujita and H. Costing, J. Amer. Chem. Soc. 79,

1099 (1956).

(b) J. S. Kirkaldy, Can. J. Phys. 37, 30 (1959).

(c) Y. Oishi, J. Chem. Phys. 43, 1611 (1965).

(d) A. G. Guy, V. Leroy and T. B. Lindemer, A.S.M. Transactions Quarterly, 59, 517 (1966).

[4] O. Hougen & K. Watson, Chemical Process Principles,

Vol. Ill, Wiley, New York, 1947, pp. 97-99.

[5] FeO Papers at this Conference: J. B. Wagner, Jr.; H. F. Rizzo, et al.; L. Himmel and J. B. Holt; R. H. Campbell M. O'Keefe.

[6] L. Onsager, Ann. N. Y. Acad. Sci., 46, 241 (1945).

[7] T. O. Ziebold, Ternary Diffusion in Copper-Silver-Gold

Alloys, Doctoral Thesis, M.I.T., Cambridge, Mass., 1965.

[8] (a) A. R. Cooper, Jr. and W. D. Kingery. J. Amer. Ceram.

Soc. 47, 37 (1964).

(b) A. R. Cooper, Jr., J. Chem. Eng. Sci., 21, 1095 (1966).

[9] V. G. Levich, Physicochemical Hydrodynamics, Prentice HaU, 1962, chapter 2. [10] A. R. Cooper, Jr. and A. K. Varshneya, to be published in

J. Amer. Ceram. Soc. 51, 103 (1968).

Only three of these are truly independent.

84
Divalent Cation Impurity Diffusion in MgO

A. J. Mortlock

Australian National University, Canberra.

1. Introduction

Direct measurements are now available of the

diffusion in MgO of Mg^+, Ni2+, Co2+, Fe-+, Zn-+

,

Ca^+, Be^+ and Ba^+ [1 to 6].^ In the single case

of Ba^+ the penetration profiles could be analysed

into two components yielding different activation

energies {Q) and pre-exponential factors [Do)-

With the exception of M^+, and Ba-+ at small

penetrations (<20/a), all the observed ^'s lie in the range 1.6 to 2.1 eV and the i)o's are roughly

of the order of 10"^ cm^/sec. The data for Mg-+,

and Ba^+ at small penetrations, yield ^'s of approximately 3.4 eV and Z)o's of the order of 10"' cm^/sec.

Harding [6] considers that the lower ^'s and

Do's are strongly influenced by dislocations.

Furthermore, in the case of Ba-+ the higher Q and

Do measured at small penetrations are associated

with lattice diffusion in this region and are only

observed because of the relatively high resolution

(~ \ Ijl) of the grinding technique used to analyse

the penetration profiles. Lidiard [7] on the other

hand has interpreted the Ni^^, Co2+, Fe-+, Ca^^

and Be^+ results as being characteristic of extrinsic diffusion and, therefore, the observed ^'s

as energies of motion, Hm- This conclusion was

reached on the basis of the maximum anneahng

temperatures used, 1700 to 1800 °C, and the purity

of the MgO crystals, nominally 99.99 percent for the Ca-+, Be^+ and Ba-+ investigations. Thus, for crystals with a concentration of ahovalent impurities of 10~* the extrinsic-intrinsic transition

is at approximately 1700 °C. The relatively high

Q for Mg2+ was explained as intrinsic diffusion

due to the unusually high purity of the crystals used

in these particular experiments.

The incompatibility of these two interpretations

of the experimental data requires closer examination. This is particularly so as the problem cannot

be easily resolved at the present time by using

higher purity MgO crystals because these are unavailable or, because of vaporization losses,

annealing at much higher temperatures. It is also

of interest to consider the reason for the unusual

' Figures in brackets indicate the literature references at the end of this paper.

two-component nature of the Ba^+ penetration

profiles.

2. Evidence for Extrinsic Diffusion

The following points throw light on the first of these problems:

(i) The MgO crystals used for the Ca2+, Be2+

and Ba-+ diffusion experiments came from the same source and were of the same nominal purity.

Therefore, as Lidiard's calculations indicate that the measured diffusion of Ca'^+ and Be^+ is in the

extrinsic region, this should also apply to the Ba-+ data, the temperature range being approximately the same in all three cases.

(ii) With the single exception of the large penetration Ba^+ results, all the available Q measure- ments can be fitted by the equation = 1.34+ (1.05 X 10i«)r2 eV (1) where r is the cation radius in A; see figure 1. For

this fitting Lidiard's [7] value of 1.5 eV for Hm for Mg^+ was used and the upward-pointing arrow at- tached to the plotted point indicates, following

Lidiard, that this could possibly be a few tenths of an eV higher. The operation of this equation demonstrates a simple variation of Hm with cross- sectional area of the diffusing cation. This in turn

indicates a consistent dependence of Hm on elastic strain energy at the saddle point.

(iii) The relatively large Do factor for the small

penetration Ba-^ measurements does not appear

to be entirely inconsistent with the much smaller

Z)o's for the other impurities when due account

is taken of the relatively large size of this cation: r= 1.35 A. This is more clearly seen if it is supposed

following Lawson [8| that S,,,, the entropy of motion,

is proportional to the activation volume of motion

and then, as a very rough approximation, that this

is a linear function of H. Under these circumstances

log Do should increase linearly with r', particularly

for the larger values of r: the experimental data

scatter badly but are not inconsistent as a whole

with the predicted tendency.

The combined evidence represented by (i) to

(iii) points very strongly to the conclusion that all the impurity cation diffusion measurements so

85
0-5 I O 1-5

,16

2-0 (ionic radius in a ^ X lO'

Figure 1. Correlation between movement energies for cation diffusion in MgO and cation radius.

far reported in MgO have been made in the extrinsic

region. This extends and further supports the similar

conclusion reached earher by Lidiard [7| on the

basis of a smaller amount of experimental data. The problem of the two component nature of

the Ba-+ penetration profiles can be explained by

noting that a similar behaviour has been observed

by Gupta and Weirick [10| for the self-diffusion of Ca in CaO. The near-surface diffusion region (0-

20 /x) is attributed in this case to extrinsic diffusion

and the deeper diffusion region to the super position

of extrinsic and dislocation diffusion. The reason why Ba-"*^ alone shows this effect in MgO could be

due to the enhanced effect of a smaller density

of dislocations, in this case, caused by the unusually

large segregation to them by the large Ba-+ ions,

Mortlock [11 1. The observed very small solubility of BaO in MgO, Harding [12|, is consistent with this.

3. Mullen's Equation

In view of the success of eq (1) in fitting the

experimental data, it is of interest to examine

the performance in this connection of more complex

equations which have been derived from strain energy considerations. Mullen [13| has shown

that the energy of motion for cation self-diffusion

in an ionic solid, H„w^ is given by

H,„o = K (l-llViyrl (2)

where K is an empirical constant and r„ is the

nearest- neighbour distance. It is also shown that

the energy of motion for an impurity cation which

is larger than the host ion it replaces is given by

H,„ = K{[in-r^) + il-llV2) r„]-'-3(r,-r^)-^}

(3)

where we now distinguish between the radius of

the impurity cation r, and that of the host cation,

/ + The second term in this equation arises out of

the gain in strain energy resulting from the relaxation of the anion sublattice which takes place when

the cation moves away from its equilibrium position.

If we put

Ar=n — r+

^H,n — H,i, — H,„o

then we can eliminate the constant K to write

AH,„ = H,„o {6.83 (^yj - 23.3 (^)'| (4)

In figure 2 the experimental data for cation

diffusion in MgO are plotted. The dashed line in this figure corresponds to the first term only in eq (4) while the full line represents both terms. This equation shows some agreement with experiment up to and including ions as large as Ca'^+ for which (Ar/ro) is 0.16. For larger ions, however, the

only data available (Ba^+) suggests that the reduction

86
•4

Figure 2. Comparison between the observed and predicted increase in energy for the movement of divalent cation impurities in MgO

relative to that for self-diffusion.

in the energy of motion due to relaxation of the

anion sub-lattice is over estimated. Overall, the

semi-empirical equation (1) seems to describe the

effect of ionic size variations on the motion energy

better than eq (4).

4. References

[1] R. Lindner and G. D. Parfitt, J. Chem. Phys. 26, 182 (1957).

[2] B. J. Wuensch and T. Vasilos, J. Chem. Phys. 36, 2917

(1962).

[3] B. J. Wuensch and T. Vasilos, J. Chem. Phys. 42, 4113

(1965).

[4] J. Rungis and A. J. Mortlock, Phil. Mag. 14, 821 (1966).

[5] B. C. Harding and A. J. Mortlock, J. Chem. Phys. 45, 2699

(1966).

[6] B. C. Harding, Phil. Mag. 16, 1039 (1967).

[7] A. B. Lidiard, Proc. British Ceramic Soc. No. 9, July, 1967,

p. 9.

[8] A. W. Lawson, J. Phys. Chem. Sohds 3, 250 (1957).

[9] R. W. Keyes, J. Chem. Phys. 29, 467 (1958).

[10] Y. P. Gupta and L. J. Weirick, J. Phys. Chem. Solids 28,

811 (1967).

[11] A. J. Morlock, Acta Met. 8, 132 (1960).

[12] B. C. Harding, M. Sc., Thesis Australian National University, Canberra (1967), p. 27.

[13] J. G. Mullen, Phys. Rev. 143 , 658 (1966).

87

Diffusion of Iron and Nickel in Magnesium Oxide Single Crystals ^

Stuart L. Blank and Joseph A. Pask

Inorganic Materials Research Division, Lawrence Radiation i Laboratory, and Department of

Mineral Technology, College of Engineering, University of California, Berkeley, California 94720

1. Literature Data

The vacancy mechanism is the most widely

accepted mechanism for diffusion among oxides.

The self-diffusion coefficient of Mg+^ in MgO was

reported by Lindner and Parfitt [1] ^ as

£) = 0.249 exp (-79.0 kcalIRT)

in the temperature range of 1400 to 1600 °C. If the

MgO was "sufficiently pure," the value for the

activation energy represents that for the intrinsic

diffusivity region. Self-diffusion measurements of

Ni"^^ in NiO by Choi and Moore [2] in the temperature range of 1000 to 1470 °C showed

£>= 1.82 X 10-3 exp (-45.6 kcal/«r)

Self-diffusion measurements of Fe+^ in Feo.gorO by

Himmel et al. [3] in the temperature range of 700

to 1000 °C showed

Z) = 0.118 exp (-29.7 kcallRT)

Diffusion measurements of iron in Q:-Fe203 in air by Izvekov et al. [4] in the temperature range 950

to 1050 °C showed

D= 1.3X106 exp (-100.2 kcaWT)

It has been found by several investigators that addition of Mg"^^ ions to FcxO results in a decrease

of the Fe^^/Fe^^ ratio indicating that magnesium

ions do not act purely as a diluent. Diffusion studies on the FcxO-MgO system, which forms a complete

solid solution, are contradictory. Wuensch and

Vasilos [5] have shown that D = 8.83 X 10-5 exp (-41.6 kcaVRT)

' Extended abstract; complete paper has been submitted for publication in the Journal of the American Ceramic Society.

^ Figures in brackets indicate the literature references at the end of this paper.

indicating no compositional dependency. Rigby

and Cutler [6] found that the diffusivity but not the

activation energy is concentration dependent.

In reaction studies between Fe203 and MgO

Carter [7] concluded that the solid-state reaction

forming MgFe204 occurred by counterdiffusion of

the Mg+^ and Fe"*^^ ions through a relatively rigid

oxygen lattice. He found the ratio of MgFe204

formed on the MgO side of the initial interface

to that formed on the Fe203 side to be 1:2.7. Fresh

and Dooling [8] indicated that the rate controlling

step in this reaction is the diffusion of the oxide

components through the ferrite product layer.

Schaefer and Brindley [9] found that "isolated

islands" of MgFe204 developed inside crystals of

(Fe, Mg)0 upon oxidation.

2. Experimental Procedure

2.1. NiO-MgO Diffusion Couples

Reagent grade NiO powder was used for all of

the polycrystaUine NiO-single crystal MgO dif- fusion runs. Single crystal NiO vs. single crystal

MgO runs were also made. Diffusion anneals were made in the temperature range 1200 to 1400 °C

both at atmospheric pressure and at 10"^ torr.

2.2. Fe-MgO Diffusion Couples

Single crystals of MgO were packed in reagent

grade iron powder and annealed at temperatures

from 1100 to 1350 °C at a pressure of 10-** torr.

2.3. Fe203-MgO Diffusion Couples

Reagent grade ferric oxide was used for all polycrystalhne Fe^Os diffusion runs. Single crystals

of MgO packed in Fe2Qi powder were annealed

in air in the temperature range 1020 to 1500 °C. A run of single crystal Fe203 versus single crystal

MgO was also made at 1330 °C. Diffusion runs were

either air-cooled or cooled with the furnace rather

290-621 0-68—

7

89
than quenched into water in order to prevent shat- tering of the MgO single cyrstals. Several samples,

however, were quenched into cold water from

1350 °C and 1500 °C for the purpose of obtaining

EPR and optical absorption specimens.

2.4. Analytical Studies

An Apphed Research Laboratories electron

microprobe was used to determine diffusion profiles.

The Ka x-ray emission line was monitored for

all the cations present. Specimen currents were

set at 0.03 fxa. The data output from the microprobe was fed to a CDC 6600 computer where the various

correction factors necessary to quantitatively convert counts to concentrations were applied.

Specimens of the Fe-MgO and Fe2 03-MgO sys- tems were prepared for the electron paramagnetic

resonance studies in the same manner as for the

diffusion studies. Crystals of MgO containing

gradients of iron were either cleaved or ground

to obtain varying concentrations of iron ions in the crystals. EPR spectra were also obtained for the MgO crystals as received and after various treatments. The EPR spectrometer was cali- brated by using two standards: Diphenyl-picryl

hydrazyl (D.P.P.H.) with a g value of 2.0037 and

5 X 10'4 spins of Cr+-^ in MgO (g= 1.987).

Optical absorption spectra were taken on speci- mens obtained from the diffusion zone of several

Fe20:i-MgO and Fe-MgO couples. A Beckman IR-4 was used for the 1-10 /a region and a Gary 14 for the 250 mfji-lfx region. Patterns were taken using

the double beam.

A FejOa single crystal versus MgO single crystal

diffusion couple was examined using the Laue

back reflection technique in order to determine

whether the resulting MgFe204 zone formed by

reaction was a single crystal. Its relative orientation was also determined.

3. Diifusion Mathematics

The Boltzmann-Matano approach was applied

for solution of the diffusion equations in the Fe-MgO

and NiO-MgO systems. A computer was used to reduce error to a mimimum in solving the equation:

The mathematics needed in order to calculate

the diffusivities in the Fe^Oa-MgO system involve

a 3-phase 2-moving boundary solution. Figure 1 shows the phase diagram and diffusion profile for this system. The theoretical solution to this problem was determined by Dorn and Blank [10].

Q- 1200

Magnesio-wustite y/^ Mognesio

Mognesio- fernte + H

/ -ferrite

emotite 7

T

1

1 /

/i

?r

1 Mognesio-wustite +

~ ' /

y ' ^ Mognesio-ferrite '1

MgO

ci

100

iFesO,

Phose mCjn 0 -4" + '

Cm,n

-Ci,o

X-0 Xji J X[[

[

Figure 1. Phase diagram and diffusion profile for three phase

diffusion problem.

The data needed for the analytical solution is not known either from this study or from the literature. An approximation thus had to be made to calculate

the D values for the interdiffusivity of Fe^'* and

Mg+-. If we assume that the original interface

between the Fe203 and the MgO is located at the

boundary j we can calculate Di (D of Fe + ' in MgO).

C| — C] () +

l-erf(y,).

erf

where yi is defined by the relation

^.i,i = y. V4A^

and all other quantities are defined on the dif- fusion profile shown in figure 1. Since for all practical purposes Ci,() = 0, then

[1-erf (y,)] = 1-erf

WWit

Knowing Ci.n, t. C^ and using the approximation

stated above that A^n.i is the initial interface, we

can calculate Di from

c,

90
The resulting diffusivities are therefore indicative

of some average composition within the diffusion zone in MgO. This approximation prevents the

determination of concentration dependence if it exists, but it is thought that it is more apphcable

in this case than the Bohzmann-Matano solution

would be to a 3-phase problem.

4. Results

4.1. Fe-MgO System

In a vacuum from 1100 to 1350 °C iron enters MgO by the redox reaction

xFe + MgO = Fe^-O + Mg f .

The relationship between In D and the concentration

of Fe. in terms of the atomic percent (a/o) of Fe,

at a given temperature was found to be linear from

2 a/o Fe to the iron concentration in MgO at the

interface which was constant for all times at a given

temperature. In this system, thus, there is no moving boundary on diffusion, and the actual and

Matano interfaces are identical. Table 1 Hsts the

diffusivity results. The above results show a trend in activation

energies which is opposite to that expected on

the basis of the energies found for the end components. 79 kcal for MgO and 30 kcal for FcjO.

The observed trend in activation energy values

prompted investigations of the NiO-MgO and

Fe203-MgO systems.

4.2. NiO-MgO Systems

Figure 2 shows a typical diffusion profile ob- tained in the NiO-MgO system using single crystals

in an air atmosphere. The Boltzmann-Matano inter150 200 250 300 350 400 450 500 550 600 650 MICRONS

Figure 2. Typical diffusion profile in the NiO-MgO system annealed in air at 1370 °C.

Table 1.

C (ah Fe) Do (cm^/sec) ^H (kcal/mole)

5 6.34 X 10-7 29.6 10 4.49 X 10-6 32.9 15 5.04 X 10-5 37.7 20 8.40 X 10-^ 43.5

Table 2.

C (Ni a/o) Do (cm^/sec) ^H * (kcal/mole)

10 6.31 X 10-"^ 42.8 20 1.05 X 10-5 42.8 30 1.96 X 10-= 42.8 40 3.16 X 10-= 42.8

face at approximately 290/u, which corresponds to the initial interface lies on the higher nickel side

of the 25 ajo concentration (the actual interface)

indicating that a flow of excess Mg^-, and an equivalent amount of 0= ions in order to maintain electro- neutrality, occurred across the actual moving

NiO-MgO interface. The results of the diffusion

study in an air atmosphere are listed in table 2. In vacuum the relation Z) = 9.82 X 10"'' exp

(—45.5 kcal//?!") was found to hold with the activation energy and diffusivities concentration independent. The activation energy values for diffusion in vacuum and in air are not significantly different

to warrant postulation of any mechanism change.

These results suggest that the concentration of

vacancies in the air experiments increases with an increase of nickel concentration due to the

formation of Ni+^. Since the activation energy was

found to be independent of composition and at- mosphere, it is postulated that no structural changes

occur in this system and that diffusion proceeds

by counter-diffusion of Ni"^^, Ni+^ and Mg+'- ions on octahedral sites. Substitutional soUd solutions are therefore to be expected.

4.3. FeaOa-MgO System

Figure 3 shows a computer-drawn typical dif- fusion profile in the Fe203-MgO system obtained

in air. The irregularities in the profile starting at a concentration of approximately 1 a/o Fe repre- sent precipitates in the crystal which probably

formed on cooUng. The discontinuous increase

in iron concentration seen at 900)U, is due to the

presence of the MgFe204 phase, and the increase near the 1000^t region is due to the transition to the Fe203 phase. Optical examination of the

91
MICRONS

Figure 3. Diffusion profile in the Fe-iO^-MgO (air) system annealed at 1460 °C for 42 hours [computer drawn).

diffusion specimens at room temperature indicated

that the precipitates in the MgO single crystal

region were coherent since they were not visible but that the precipitates of a:-Fe203 in the MgFe204

phase were incoherent.

The x-ray analysis of a single crystal versus

single crystal couple showed that the MgFe204

phase formed by reaction was a single crystal

oriented with the (100) face of the MgFe204 coincident with the (100) face of the MgO. The observation with the electron microprobe that several

Fe203 precipitates were in contact with the MgO

region without reaction substantiated the conclusion that the precipitates in the MgFe204 region were formed on cooling. A large solubility of Fe203, probably as Fe304, in the MgFe204 at high

temperatures is therefore indicated. The data also showed that the solubility of MgO or MgFe204 in Q!-Fe203 is very small. Because the limit of de- tection of precipitates by the electron probe was

on the order of the precipitates were of that order of size in the 1 ajo Fe concentration region.

The results of the diffusion analysis showed the

relationship

D = 8.83 exp (- 74.6 kcal//? T)

to hold over the temperature range 1000 to 1460 °C.

This expression is the mathematical approximation

for some average composition within the diffusion zone in MgO. The analysis did not allow a deter- mination of any concentration dependence of either

the Do or the activation energy terms if it existed.

4.4. Electron Paramagnetic Resonance

Studies

Examination of the MgO starting crystals by

EPR showed that octahedraUy coordinated Fe+^

ions were present. These could be reduced to Fe^^

by heating in vacuum (10~*' torr) for 48 hours at 1200 °C. The spectroscopic analysis showed a

total iron concentration of 0.01 percent. The

^--value for Fe+^ at the 0.01 percent total iron concentration was found to be 1.993. Further studies

by EPR on crystals to which small amounts of iron were added by diffusion showed that the g value

for the Fe+^ ions moves from 1.993 (V2—> — V2

transition) to higher g values as the trivalent iron

concentration increases. These studies also showed

that at 0.2 a/o Fe^^ and below it is relatively easy

to reduce all the Fe+^ ions to Fe+^ by heating in a vacuum. Specimens were also investigated from the dif- fusion zone in MgO of the Fe203-MgO couples.

The EPR spectra of a sample from a couple run

in air and containing a high iron concentration

shows two important features: a sharp transition

occurring at ^=2.56 and a broad absorption

starting at approximately 1000 gauss. The spectra

of the same specimen after reduction at 1300 °C

for 10 hours in vacuum (10"^ torr) shows that the

sharp transition at ^=2.56 disappears, and that the broad absorption remains but is shifted to the

g=2A^ region. In both cases the specimen was

ferrimagnetic, indicating that some Fe+^ ions were

in tetrahedral sites. From the spectra discussed

and other EPR experiments the following assignments were made for the important features ob- served in the patterns. The peak at^= 1.993 which

broadens and shifts to ^=2.56 is the V2^— V2

transition of Fe^^ coordinated in a distorted octa- hedral site. The ^--value shift is due to the ferri- magnetic MgFe203 elements which form at the

higher iron concentrations. The field as seen by

an Fe"*"^ ion in such an environment will be considerably different from the field seen by an Fe+^

ion in a paramagnetic environment. The ferri- magnetic centers tend to concentrate the field and

therefore cause the ¥2^ — V2 transition to occur

at a lower externally applied field. The broad

absorption starting about 1000 G is related to the

ferrimagnetic material in the crystals.

EPR spectra in the Fe-MgO system showed

similar features. Both octahedral Fe+^ and ferri- magnetic material were present.

4.5. Optical Studies

The optical investigation showed that Fe+^

(octahedrally coordinated) has a broad absorption

band centered at 1/u,. Trivalent iron shows a band

in the region of 0.5/Lt. Specimens from an Fe203-

MgO couple quenched from 1350 °C in air showed

a different composition when compared with similar

specimens that were slow-cooled indicating that compositions at the annealing temperatures are

92
not necessarily the same as those measured at room temperature. The quenched samples were

higher in Fe+^ content than those that were slowcooled.

5. Discussion

The results of the Fe-MgO diffusion experiments

under vacuum conditions indicated that both Do

and the activation energy vary with iron concentration. Of particular interest are the values for the activation energy which increase from 29 kcal

at 5 alo to 43 kcal at 20 a/o Fe. The EPR and

optical studies on specimens cooled in vacuum

have shown that the introduction of Fe+- ions into MgO is associated with the appearance of Fe+^

ions by reaction, with some unknown amount in tetrahedral sites. The concentration of thermally created vacancies

in the MgO was compared with that chemically

created due to the presence of original trivalent impurities. Assuming that the heat of formation of

a Shottky defect in MgO is 5 eV, the equihbrium

concentration of vacancies in theoretically pure

MgO at 1200 °C would be about lO'^/cm^. If the

only trivalent impurities in our crystals were the

0.005 wt percent AI2O3 indicated by spectroscopic

analysis, then the calculated concentration of

chemically created vacancies is a minimum of

lO'Vcm^, or 4 orders of magnitude greater than

the equilibrium concentration of the thermally

created vacancies. The existing impurities therefore place all diffusion experiments into the extrinsic rather than intrinsic region. This result then provides an explanation for the A// value of

29.6 kcal/mole at the low iron concentrations rather

than the expected 79 kcal for the intrinsic region.

Examination of the MgO structure indicates that the four octahedral sites adjoining a tetrahedral

site filled with an Fe^^ ion must be vacant in order

to avoid extremely high cation-cation repulsive

forces. The resulting defect in the MgO matrix is now an embryonic element of the inverse spinel structure whose size initially would be less than one unit cell of the MgF6204 structure. These

embryonic units, however, could act as nucleation

sites for subsequent precipitation in the system

on cooUng. It would be expected that the bond

strength of the Mg+'^ ions coordinated with oxygens

associated with such fiUed tetrahedral sites would

be significantly increased. This would result in an

increased heat of motion term when these Mg+^

ions move ' since counterdififusion must occur to

maintain electroneutrahty. This structural change

occurring inside the crystal could then provide an

explanation for the increasing activation energy

with increase of iron seen in this system.

The results obtained in the NiO-MgO system,

however, suggest a different mechanism. In this case the dififusivity but not the activation energy

is a function of the nickel concentration in air; in vacuum there is no concentration dependence.

These results and the absence of precipitates in the system suggest that no structural changes occur and that most, if not all, of the cations remain on

octahedral sites. In this case a true substitutional

solid-solution exists. The optical studies have shown the existence of Fe+- in the FeaOa-MgO system. The mechanism for formation of Fe+- from the Fe+^ ions is beheved to be related to the formation of MgFe204 in the sys- tem. In this process a driving force is estabhshed

for the reduction of some Fe+^ to Fe"*^- with the

release of oxygen at the surface. The oxidationreduction process should be relatively easy due to the small energy necessary to transfer an electron

between neighboring octahedrally coordinated Fe+^

and Fe+^ ions. This mechanism could also be used

to explain the effect observed with addition of

Mg+2 ions on the Fe+3/Fe+2 ratio in magnesiowiistite crystals.

6. References

[1] K. Lindner and G. D. Parfitt, Diffusion of Radioactive Magnesium in Magnesium Oxide Crystals, J. Chem.

Phys. 26, 182 (1957).

[2] J. S. Choi and W. J. Moore, Diffusion of Nickel in Single

Crystals of NiO, J. Phys. Chem. 66, 1308 (1962).

[3] L. Himmel, R. F. Mehl and C. E. Birchenall, Self-Diffusion

of Iron in Iron Oxides and the Wagner Theory of Oxidation, Trans. Am. Inst. Min. Met. Eng. 197, 827 (1953).

[4] V. I. Izvekov, N. S. Gorbunov and A. A. Babed Zakhryapin,

Diffusion of Iron in Hematite, Fiz. Metal, i MetaUoved

14, 195 (1962).

[5] B. J. Wuensch and T. Vasilos, Diffusion of Transition Metal Ions in Single-Crystal MgO, J. Chem. Phys. 36,

2907 (1962).

[6] E. B. Rigby and I. B. Cutler, Interdiffusion Studies of Fe^O-MgO, J. Am. Ceram. Soc. 48, 95 (1965).

[7] R. E. Carter, Mechanism of Sohd-State Reaction Between Magnesium Oxide and Aluminum Oxide and Between Magnesium-Oxide and Ferric Oxide, J. Am. Ceram.

Soc. 44, 116 (1961).

[8] D. L. Fresh and J. S. Dooling, Kinetics of the Solid-State Reaction Between Magnesium Oxide and Ferric Oxide,

J. Phys. Chem. 70 , 3198 (1966).

[9] W. L. Schaefer and G. W. Brindley, Oxidation of Magnesiowiistite Single Crystals, J. Phys. Chem. Solids 24, 919

(1963).

LlO] J. Dorn and S. L. Blank (to be published).

93

Impurity Cation Diffusion in Magnesium Oxide

B. J. Wuensch

Ceramics Division, Department of Metallurgy and Materials Science

Massachusetts Institute of Technology, Cambridge, Massachusetts 02139

and

T. Vasilos

Avco Corporation, Space Systems Division

Lowell, Massachusetts 01887

Diffusion coefficients for Ca.^+ in single-crystal MgO have been determined in the range 790 to 1850 °C from concentration gradients obtained with the aid of electron microprobe spectroscopy.

Step-function distributions of solute provided artificial Gaussian gradients with an apparent Z^f product

of 1.7 10~* cm^, and demonstrated that the microprobe technique could not be reliably used for the

analysis of gradients less than lOju, in extent. A procedure for correction of the gradients for this effect was developed, and will be applied to the data. The diffusion coefficients obtained are tentatively represented by an activation energy of 2.76 eV, and a pre-exponential term D„ of 8.9 10"'' cm^/sec.

Diffusion coefficients obtained above 1200 °C are in good agreement with recent results of Rungis and Mortlock. Data are available for both anion and cation self-diffusion in MgO, and for the diffusion of S_dit

ferent impurity cations. None of these results have revealed a change of slope in the temperature

variation of the diffusion coefficient. Unequivocal identification of the diffusion processes as extrinsic or intrinsic is difficult. Available estimates of the energy required for Schottky pair formation in MgO

indicate the purity of the crystals employed in most studies should result in extrinsic behavior at all temperatures up to the melting point of the material. This view is supported by activation energies

for impurity ion diffusion which are less than half the energy required for defect formation, and small Do values characteristic of extrinsic diffusion. The Mg^"^ self-diffusion data appears to represent in- trinsic transport. The energy for Mg^+ self-diffusion should therefore display no correlation with

energies for impurity ion diffusion or for cation self-diffusion in transition metal oxides. Recent cation

diffusion data do not display a regular variation with the ratio of ionic radius to polarizability, as pre- viously proposed by the writers. It is concluded that the early agreement with this relationship was

fortuitous. Interpretation of the variation of impurity cation activation energy in terms of a size effect explains the difference in values observed for ions of greatly different size, but cannot explain smaller

differences between ions of similar radius. It is proposed that different correlations may apply to spherical-shell ions and to ions, such as the transition metals, with highly polarizable d electrons.

Wuensch & Vasilos:

Key Words: Cation diffusion, electron microprobe, magnesium oxide, Schottky pair.

1. Introduction

Many properties of ceramic oxides are controlled

by the rates of diffusion of one of the constituent

ions. Both anion and cation self-diffusion rates have

therefore been determined for a number of materials in order to gain insight into the kinetics of

such processes. Study of impurity ion diffusion, on the other hand, affords the possibility of deter- mining the influence of the nature of the diffusing

atom on diffusion in a given oxide when such meas- urements are made for a series of different ions. Correlations between diffusion rates and properties

of the migrating species have been observed for metals. No completely satisfactory relationship

has been established for ionic crystals.

Diffusion rates have now been determined for a number of impurity cations in single-crystal MgO.

This oxide has been studied for a number of reasons. Magnesium oxide has the NaCl-type

structure common to many of the alkali hahdes.

Transport has been extensively studied in these

materials and is relatively well understood. Single

crystals of MgO, although of questionable character,

are readily available. Oxides of a number of divalent

cations display complete or appreciable solid solu95
bility in MgO, and thus are amenable to study. The

present paper describes preliminary results of a study of diffusion of Ca-+ in single-crystal MgO,

and summarizes the data presently available for the

diffusion of other impurity ions in this material.

2. Experimental

2.1. Sample Preparation

Diffusion couples in the system CaO-MgO were

prepared with single crystals obtained from the

Norton Company, Worcester, Mass. Specimens

of two varieties were employed. In a first type, a

1/u, layer of 50 wt percent CaO-MgO solid solution was placed between two polished (100) surfaces

of MgO single crystals. The vapor pressure of CaO

was appreciable at the higher temperatures investigated, and it was necessary to encapsulate the

entire specimen in a cylinder of pressed MgO

powder. This compact sintered to high density

during the early portion of the diffusion annealing

and prevented loss of solute through vaporization.

Annealing of the thin-film type of diffusion couple

produces a Gaussian distribution of solute. If the

diffusion coefficient, Z), is independent of concentration, a plot of the logarithm of solute concentration

as a function of the square of penetration is linear,

and has slope given by (4£)f)~', where t is the time

of the diffusion anneahng.

In a second type of specimen a poHshed (100)

surface of a MgO crystal was placed over a small

well of sintered MgO in which a pellet of 50 wt

percent CaO-MgO solid solution had been placed.

The surface of the crystal was thus in equihbrium

with a vapor of the solute and was maintained at constant concentration throughout the course of the anneahng. If the diffusion coefficient is independent of concentration, these boundary conditions produce a solute distribution given by Co

erfc :t;(4Z)0"''", where Co is the surface concentration and erfc is the complementary Gaussian error

function. The inverse complementary error function

of CjCo is a linear function of penetration into the

sample, and has a slope given by (4/>f)~'/^. All diffusion anneahngs were conducted in air. Samples were prepared in the temperature range

790 to 1850 °C. The duration of the annealings

ranged from 24 hr, at the higher temperatures

investigated, to as long as 1.2 yr.

surface was then pohshed. The distribution of

CaO solute was determined with the aid of elec- tron microbeam probe spectroscopy. The samples were analyzed with a 30 kV electron beam having

a diameter of approximately l)Lt. Intensities of

CaKa radiation were recorded relative to the

intensity obtained from a CaCOs standard to ehminate possible variations in beam characteristics.

The intensity ratios were converted to concentrations of CaO through comparison with intensities obtained from a set of fine-grained sintered CaOMgO standards.

Diffusion rates at the lower temperatures which were investigated were extremely small. Diffusion

anneahngs in excess of a year in duration produced

measurable concentration gradients of only 10 to

15/A in extent. Since the diffusion coefficients ob- tained from the initial samples prepared at these

temperatures were beUeved to be less accurate than those obtained at the higher temperatures,

additional samples were prepared which employed

diffusion anneahngs of longer duration. The latter samples usually provided diffusion coefficients

which were smaller than those obtained from comparable specimens annealed for shorter times.

It was therefore suspected that the short gradi- ents which had been measured had been influenced

by lack of resolution of the microprobe. This was

experimentally confirmed by coating MgO crystals

with a thin layer of solute in the manner used to prepare diffusion couples. These specimens were

not subjected to a diffusion annealing. Normal

microprobe analysis of crystals with such a stepfunction distribution of solute provided an ap- parent concentration gradient approximately 8/a

in extent. The apparent solute distribution was

Gaussian in nature, and thus was difficult to dis- tinguish from actual solute penetration into the

material. The form of the apparent distribution was independent of the concentration of CaO in the step function, and corresponded to an apparent

product Dt of 1.7 lO"* cm^.

This effect provides an example of the caution

which must be exercised in examining small-scale

features through microprobe spectroscopy. In matrices of low atomic number, such as MgO,

penetration of the electron beam into the specimen

and secondary fluorescence may decrease the ef- fective resolution of the electron beam to several

times the nominal diameter of the electron beam.

2.2. Diffusion Couple Analysis

After completion of the diffusion anneahng,

the MgO crystals were cleaved along a (100) plane

parallel to the diffusion direction. The cleaved

2.3. Correction of the Data

The concentration gradients measured with the

microprobe correspond to the convolution of the

true concentration gradient with an instrumental

96
broadening function which was approximately

Gaussian in nature. The form of the apparent concentration gradient obtained upon analysis of stepfunction distributions of solute was reproducible

and independent of concentration in the step. It was therefore assumed to represent a reliable measure of the instrumental broadening.

Correction for the effect is simple for thin film diffusion couples in which all solute concentrations

fall within the permissible range of sohd solution

of CaO in MgO. The convolution of two Gaussian

functions is a Gaussian function with an exponential

parameter equal to the sum of the exponential

parameters of the two original functions. A true value of the product Dt may therefore be obtained

by subtracting from the apparent value the product

1.7 10~^ cm^ obtained from the step functions. This

procedure unfortunately is not valid for the complementary error function distributions, or for several of the thin-film specimens which displayed a surface layer of CaO which had not been completely

dissolved in the MgO crystal.

For these specimens it was assumed that each

apparent concentration, C", experimentally deter- mined at a discrete penetration Xi, was given by

the sum of the true concentration, C", plus a linear

combination of the true concentrations at neighboring penetrations. This relation may be expressed

C"(xO=C»U.) + 2^jC''U>j)

where the proportionahty constants, k, are deter- mined from the apparent concentration distribution

obtained from the step function. This provides a set of simultaneous equations which may be solved

for the unknown concentrations C". Details of this procedure, and comparison of the corrected and

apparent concentration gradients will be described

elsewhere [1].'

3. Results

Figure 1 presents a concentration distribution

typical of those obtained from analysis of the vapor- deposition specimens prepared at higher temperatures. Surface concentrations ranged between 0.1 •and 0.9 atomic percent CaO. The inverse complementary error function of CjCo varies linearly with

' Figures in bracltets indicate the literature references at the end of this paper.

^ It might be expected that the correction procedure would provide diffusion coefficients which would be still smaller, since the apparent gradient is larger in extent than the true distribution of solute. Most of the low-temperature specimens, however,

retained a layer of undissolved CaO at the surface of the crystal because of the reduced

range of solid solubility at these temperatures. The true concentration gradient is therefore masked by a steep apparent gradient contributed by this layer. The cor- rected gradient, while smaller in extent, has a much smaller slope. The corrected

diffusion coefficient is therefore larger.

penetration into the specimen indicating that,

within experimental error, D is not a function of

concentration. A representative concentration dis- tribution for the thin-film type of specimen is given

in figure 2. The corrections for instrumental broadening of the distributions were negligible for the more extensive penetrations obtained at high temperatures. The diffusion coefficients obtained are

plotted as a function of reciprocal temperature in figure 3. Rungis and Mortlock [2] recently performed a similar set of measurements using MgO crystals

obtained from a different source. The techniques

employed in their study were quite different than

those of the present work. Diffusion of the radio- active tracer Ca^"' was studied, and concentration

distributions were measured through autoradiography. Despite these differences in approach, the

results of the two studies are in reasonable agree- ment in the temperature range 1200 to 18.50 °C. For lower temperatures, however, the present work

provides diffusion coefficients which are considerably smaller. The results of the present study are therefore described by

D = 8.9 lO--" exp (-2.76 eV/AT)

while Rungis and Mortlock's data are represented by

I> = 2.95 10-^ exp (-2.13 eV/AT).

Our low temperature data, however, are based on

concentration gradients not yet corrected for in- strumental broadening. The correction will increase

these diffusion coefficients and bring the results of

the two studies into closer accord.

4. Discussion

4.1. Comparison of Diffusion Parameters

Both cation [3] and anion [4] self-diffusion rates have been determined for single-crystal MgO. Data

are also available for the diffusion of a number of

impurity ions. These include Ni-+ [5], Co-+ [5],

Fe^+ [5, 6], Mn^+ [6], Zn^+ [7], Be^+ [8], and

Cr''+ [6]. The properties of these ions, and the

parameters describing their diffusion in MgO are summarized in table 1. Not included in this compilation are the results of several additional studies

which have employed diffusion couples of different geometry. In these experiments a single crystal of MgO was placed in contact with an extended source

of polycrystalhne solute oxide. These samples pro- vide coefficients for interdiffusion between MgO

and a pure solute oxide which, being polycrystalline,

is more difficult to characterize. Measurements of

97
PENETRATION (MICRONS)

Figure 1. Plot of the inverse complementary error function of CaO concentration divided by concentration at the sample surface as a function ofpenetration into single crystal MgO for a specimen with constant surface-concentration boundary conditions (Co expressed

as atomic fraction CaO).

PENETRATION (MICRONS)

100 150 200

10 20 30 40 50 PENETRATION SQUARED (MICRONS SQUARED Xio')

Figure 2. Plot of the logarithm of concentration of CaO as a function of the square of penetration into single-crystal MgO for a specimen with thin-film boundary conditions.

this type have been conducted for Ni^+ [9-11],

Co^+ [9]. Fe^+ [11, 12]. and Mn^+ [la]. The principal difference between the results obtained from

this type of specimen and those obtained from the

thin-film or vapor-deposition type is that the dif- fusion coefficients obtained from the former have

been found to be a marked function of concentration. None of the experiments with vapor-deposition

or thin-film diffusion couples [2. 5-8]. have pro- vided evidence for a concentration-dependent dif- fusion coefficient. The explanation for this apparent

discrepancy may be in the range of concentrations

explored with the differing types of specimens.

The work summarized in table 1 has been concerned

with low solute concentrations. In previous work

by the writers, for example, surface concentrations

of solute oxide rarely exceeded 10 atomic percent.

The principal portion of the concentration profiles

consisted of concentrations of the order of frac- tions of an atomic percent.

In early work [5] the writers noted that the activa- tion energies for diffusion of transition metal ions

in single-crystal MgO decreased as the radius of

the migrating ion increased. To explain this be- havior, it was assumed that the contributions to the barrier to diffusion which differed most for

98
TEMPERATURE (DEGREES CENTIGRADE)

800 900 1000 1100 1200 1300 1400 1600 1800 2000

10.0 90 8.0 70 6 0 5.0 4 0

temperature' (DEGREES KELVIN ' X 10*

) Figure 3. Plot of diffusion coefficients for Ca '^* in single-crystal MgO as a function of reciprocal temperature.

Triangles represent diffusion coefficients obtained from constant surface-concentration specimens, circles values obtained from thin-film samples.

Table 1. Data for diffusion in single crystal MgO

Ion lonic radius, r Polarizability, a rja E Do Material Reference

(10-" cm) dO--^" cm^) (10'" cm-2) (eV) (cm^/sec)

.31 .012 25.83 1.60 1.41 10-' M Harding & Mortlock [8| Cr'* .63 .852 .74 2.95 9.8 10-^ N Tagai, et al. [6| Mg^^ .65 .111 5.86 3.43 2.49 10-' N. I Linder & Parfitt [3| Ni2+ .69 .246 2.80 2.10 1.80 10--' N Wuensch & Vasilos [5| Co^+ .72 .289 2.49 2.06 5.78 10-^ N Wuensch & Vasilos [5| Zn2+ .74 .250 2.96 1.85 1.48 10-5 N Wuensch & Vasilos [7|

.75 .456 1.64

rl.8l [1.82

8.83 10-'

3.2 10-^ N

N

"Wuensch & Vasilos [5| Tagai, et al. [6| Mn-+ .80 .543 1.47 1.21 4.1 10-' N Tagai, et al. [6|

.99 .551 1.80 f2.13 12.76

2.95 10-'

8.9 10-" MN

Rungis & Mortlock [2| Present work

Co- 1.40 2.63 .53 2.71 2.5 10-" Oishi & Kingery [4|

1 = Infra Red DevelopmenI Co.. Welwyn (iarden City. Ertfiland M= Mnnucrystals (^o.. Cleveland. Ohio N = Norton Co.. Worcester, Massachusetts

99
the series of ions were the repulsive overlap energy

between the jumping ion and its neighbors, and the

polarization of the migrating ion. Increased radius

of the diffusing ion should increase the contribution

of the first term and increase the barrier to diffusion,

while increased polarizability should decrease the

barrier. When the logarithm of the activation energy

for diffusion was plotted as a function of the ratio of ionic radius, r, to polarizability, a, for the diffusing

ion, a linear relation was found for Ni'^+, Co^+,

Fe^+, and Mg^+. A similar relation [17] figure 4, exists between r/a and the activation energies for cation self-diffusion in the oxides MgO [3], NiO [14,

15], CoO [16], CaO [17], and FeO [16, 18]. Later

data [2, 7] obtained for diffusion of the impurity

ions Zn^+ and Ca'^+ provided less convincing correlation. As shown in figure 4, the data for Mn^+

diffusion [6], Be-+ diffusion [8], and Ba-+ diffusion

[19] in MgO fail completely to fit into this scheme.

It must be concluded that the correlation observed

between activation energy and r/a for the limited

amount of early data was fortuitous, and this parameter is of questionable value in establishing relations

between activation energies.

Assumption that the activation for diffusion

should increase as a function of ionic radius alone

satisfactorily explains the large difference in activation energies for ions of greatly dissimilar radius

(e.g., Be-+ and Ba-+). This cannot account, however,

for the smaller (but apparently significant) uniform

decrease of activation with increasing radius in a

series of ions of similar radius (e.g., Ni^+, Co^"*^, Zn^+, and Fe-'+). It will probably be necessary to consider additional contributions to the activation

energy for diffusion in order to satisfactorily explain

these variations. Changes in activation energy for spherical-shell ions (e.g., Be-+, Zn-+, Ca^+, and

Ba''+) might be explained solely in terms of a size

effect. Polarization of the d electrons of the transi- tion metal ions, however, will provide an additional

contribution to the activation energy for diffusion

which could easily be of the order of a few tenths

of an electron volt. Such ions might therefore

depart significantly from correlations among data

for spherical-shell ions.

4.2» Nature of the Diffusion Process

Interpretation of the mechanism of diffusion is compHcated by the fact that none of the studies

performed to date have revealed ranges of temperature within which diffusion could be described by

different activation energies. This has prevented

unambiguous identification of the activation energy

in any one diffusion study as representing intrinsic or extrinsic behavior.

A theoretical estimate of the energy required

for the formation of Schottky defects in MgO [20] has provided a value of 5±1 eV. On the basis of

this value one would predict that a concentration

of aliovalent impurity of the order of 200 ppm would

be sufficient to cause transport in MgO to be extrinsic for all temperatures up to the melting point

of the material. Most of the diffusion studies per- formed with MgO, table 1, have employed crystals

obtained from the Norton Company. Several independent analyses of this material have shown that these crystals may contain from 0.06 to 0.3 percent

impurity oxides [6, 21]. There would be little possibihty of observing intrinsic diffusion in such

crystals. The low values (ca. lO^^cm^/sec) obtained

for Do for the diffusion of all impurity ions and for 0^~ is also indicative of extrinsic diffusion. All activation energies are of the order of, or less than

half the energy for Schottky pair formation. This

would preclude possibility of intrinsic diffusion

via a Schottky defect mechanism.

In contrast, the data for Mg^+ self-diffusion are represented by a Do characteristic of intrinsic dif- fusion. Furthermore, the activation energy for diffusion may be satisfactorily interpreted as the sum of half the estimated energy for Schottky defect

formation (2 to 2.5 eV) plus a reasonable energy

(1.4 to 0.9 eV) energy for ion migration. The data

for Mg^+ self-diffusion, unlike all other studies of

diffusion in MgO, therefore appear to represent

intrinsic diffusion. The Mg^+ self-diffusion data

should therefore display no correlation with either

the extrinsic data obtained for impurity ion dif- fusion or that obtained for cation self-diffusion in transition metal oxides. (In the latter oxides it is known that stoichiometry controls the defect concentration, and the activation energy for diffusion

represents the energy for ion migration.) This

provides further evidence that the rja correlation

must be fortuitous. If the bulk of the available data

indeed represent extrinsic diffusion, there is no

basis for comparison of Do values. Such results

will primarily reflect the impurity content of the

particular crystals employed.

Problems still remain in relating the Mg^+ self- diffusion data with the results for impurity ion dif- fusion. The activation energies for impurity ion

diffusion, if assumed to represent the energy for migration of the ion, are rather high compared to that estimated for Mg'+. The available Mg-+ self- diffusion data were measured with crystals ob- tained from two different sources [3]. The manufacturer of one set of crystals stated purity of a level

for which intrinsic behavior could be expected.

The other crystals which were employed were ob- tained from the Norton Company. No difference

between diffusion rates in the two materials was

100
noted, as should be expected for intrinsic behavior.

Curiously, however, all other investigations con- ducted with the Norton crystals appear to have pro- duced extrinsic behavior.

Cation self-diffusion measurements are extremely

difficult to perform with MgO because the longestlived radioisotope of Mg has a half-life of only

21.3 hr. This has restricted measurement of dif- fusion coefficients to a narrow range of temperatures

(1400 to 1600 °C) for which appreciable tracer penetrations could be obtained within a short period

of time. The relation between the Mg^+ self-diffusion data and the data now available for impurity ion

diffusion would be considerably clarified if the

self-diffusion measurements could be extended to lower temperatures to reveal extrinsic behavior or, alternatively, impurity ion diffusion measurements

with crystals of higher quality could be extended to high temperatures to reveal an intrinsic region of

behavior.

The writers are pleased to acknowledge the

assistance of W. R. Mitchell in the preparation

of samples. This work was supported, in part, by

contract Nonr-4267(00) with the Office of Naval

Research.

5. References

[1] B. J. Wuensch and T. Vasilos, Diffusion of Ca^* in Single

Crystal MgO (to be published).

[2] J. Rungis and A. J. Mortlock, The Diffusion of Calcium in Magnesium Oxide, Phil. Mag. 14, 821-827 (1966).

[3] Roland Lindner and Geoffrey D. Parfitt, Diffusion of Radioactive Magnesium in Magnesium Oxide Crystals, J. Chem.

Phys. 26, 182-185 (1957).

[4] Y. Oishi and W. D. Kingery, Oxygen Diffusion in Periclase

Crystals, J. Chem. Phys. 33, 905-906 (1960).

[5] B. J. Wuensch and T. Vasilos, Diffusion of Transition Metal Ions in Single-Crystal MgO, J. Chem. Phys. 36,

2917-2922 (1962).

[6] H. Tagai, S. Iwai, T. Iseki and M. Saho, Diffusion of Iron, Manganese and Chromium Oxides into Single Crystal

Magnesia, Radex Rundschau 4, 577-583 (1965).

[7] B. J. Wuensch and T. Vasilos, Diffusion of Zn^+ in Singlecrystal MgO, J. Chem. Phys. 42 , 4113-4115 (1965).

[8] B. C. Harding and A. J. Mortlock, Diffusion of Be in MgO,

J. Chem. Phys. 45 , 2699-2700 (1966).

[9] I. Zaplatynsky, Diffusion of Co^+ and Ni^^ in Magnesium

Oxide, J. Am. Ceram. Soc. 45 , 28-31 (1962).

[10] J. S. Choi, Diffusion of Nickel in Nickel Oxide and Magnesium Oxide, Ph. D. thesis, Indiana University (1963).

[11] Stuart L. Blank and Joseph A. Pask, Diffusion of Iron and

Nickel in Magnesium Oxide Single Crystals, These

Proceedings.

[12] E. Bertrand Rigby and Ivan B. Cutler, Interdiffusion Studies

of the System FexO-MgO, J. Am. Ceram. Soc. 48, 95-99

(1965). [13] John T. Jones and Ivan B. Cutler, Interdiffusion Studies in the System Mn^O-MgO, Bull. Am. Ceram. Soc. 45 , 363

(1966). [14] Moon Talk Shim and Walter J. Moore, Diffusion of Nickel

in Nickel Oxide, J. Chem. Phys. 26, 802-804 (1957).

[15] Jae Shi Choi and Walter J. Moore, Diffusion of Nickel in Single-crystals of Nickel Oxide, J. Phys. Chem. 66,

1308-1311 (1962).

[16] R. E. Carter and F. D. Richardson, An Examination of the Decrease of Surface-activity Method of Measuring Self- diffusion Coefficients in Wustite and Cobaltous Oxide,

J. Metals 6, 1244-1257 (1954).

101
[17] Y. P. Gupta and L. J. Weirick, Self-diffusion of Calcium

in Single Crystal Calcium Oxide, J. Phys. Chem. Solids 28, 811-821 (1967).

[18] L. Himmel, R. F. Mehl, and C. E. Birchenall, Self-diffusion of Iron in Iron Oxides and the Wagner Theory of Oxidation, J. Metals 5, 827-843 (1953).

[19] A. J. Mortlock, Divalent Cation Impurity Diffusion in MgO, These proceedings.

[20] Jiro Yamashita and Tatumi Kurosawa, Formation Energy

of Lattice Defect in Simple Oxide Crystals, J. Phys. Soc. Japan 9 , 944-953 (1954).

[21] R. J. Stokes, Thermal-mechanical History and the Strength

of Magnesium Oxide Single Crystals: I, Mechanical

Tests, J. Am. Ceram. Soc, 48, 60-67 (1965).

102
Cavity Formation in Magnesium Oxide

A. Briggs and D. H. Bowen

Ceramics Division, Harwell, Didcot, BerksCavities have been produced in single crystals of MgO by two different processes; by annealing

certain crystals in a reducing atmosphere and by annealing crystals irradiated with fast neutrons to doses in excess of 10^" nvt. Cavities produced by the first process are lenticular in shape with diameters up to ~ 100/Lim, and contain hydrogen at pressures up to 250 atm, even when annealing is conducted

in a hydrogen-free reducing atmosphere. Those produced by the second process are cuboidal, with

sides ~ 200 A long, and contain helium and neon produced by transmutation. Preliminary studies of the cavities by a number of techniques are described and their modes of formation and the behaviour of the associated gases are discussed.

Key Words: Annealing, cavities, diffusion, magnesium oxide.

1. Introduction

The formation of a cavity in a crystalline solid requires an adequate supply of vacancies. The

vacancies may result either from lattice displacements produced by ionising or heavy particle radia- tion or from departures from stoichiometry, as for example, in crystals containing impurity ions of

different valence from those of the host lattice. These processes may be accompanied by the formation of a gaseous phase by conversion of the anion

constituent to free gas, decomposition of impurities

or by transmutation. Studies of the mechanisms of

cavity formation and the behaviour of the associated

gases may prove valuable as a means of gaining

information on transport processes in solids. Cavities are known to form in alkah hahdes under

a variety of treatments such as X and electron

irradiation, (Hibi et al. [1],' Tanaka et al. [2]), neutron irradiation, (Gilman and Johnston [3]) and an- neaUng in hydrogen, (Amelinckx et al. [4]) but apart

from studies of gas bubble formation, (Whapham

[5]), (Woollaston and Wilks [6]), comparable phe- nomena do not appear to have been reported

hitherto in oxides. This paper described the formation of cavities in single crystals of magnesium

oxide by two methods; by anneahng in a reducing

atmosphere and by annealing after irradiation with

fast neutrons to doses in excess of 10-" nvt (2=1 MeV).

In both cases the resulting cavities contain gases.

2. Cavities Produced by Annealing

in a Reducing Atmosphere

Certain MgO crystals, after annealing for periods

of several hours in a reducing atmosphere at tem-

' Figures in brackets indicate the literature references at the end of this paper.

peratures above 1000 °C have been found to contain

cavities ~ lOO/xm in diameter. These he parallel

to {100} planes, are lenticular in cross-section and,

in the plane of the cavity, are bounded by (110)

directions. (Figs. 1 and 2.) The figures also show

that the crystal surrounding a cavity is heavily

dislocated and contains a high concentration of

impurities, (Bowen [7]). Smaller cavities, ~ 1000 A

in size, are also observed that initially are octahedral

with {111} facets. (Fig. 3.) At a slightly later stage

in growth, these cavities may also develop {100}

facets. Figure 4, showing different cavities in various

stages of formation observed by ultramicroscopy,

indicates a possible sequence for growth of the

lenticular cavities. Figure 4 (1) is a region of crystal

having roughly the shape and dimensions of the

final cavity but consisting of a large number of

small cavities that act as scattering centres and

hence show as bright spots. Formation of the final cavity then appears to take place by amalgamation

of the small cavities. The amalgamation (dark area) occurs preferentially at the centre of the region

(fig. 4(3)) or sometimes in lobes directed along

either (110) or (100). Figure 5 is a scanning

electron micrograph of an embryonic lenticular

cavity cleaved through its mid-plane. In this case, amalgamation has occurred in lobes along ( 100

)

and stereo photographs confirm that the dark areas

a, b, c, d, are depressions while the bright areas

are prominences, i.e., the cruciform structure

consists of small cavities separated by walls of

crystal.

Bubbles of gas are liberated from cavities in both slowly and rapidly cooled crystals when dis- solved in acid and comparison of the sizes of

cavities in the sohd and corresponding bubbles in

103
(a) (b) (c)

Figure 1. a. Crystal cleaved on { 100} plane to reveal cross-section of a lenticular cavity.

b. As (a ) after etching to reveal dislocations.

c. Matched cleavage face after etching to reveal impurities.

the acid shows that the gas pressure in a cavity at room temperature is ~ 250 atm. Analysis, by mass

spectrometer, of the gas released when the crystals are crushed shows that it is invariably hydrogen

even when the annealing atmosphere is hydrogenfree deuterium or carbon monoxide. Not all MgO

crystals yield cavities when annealed, but those

that do exhibit a relatively large OH" absorption

in the infrared which decreases considerably on

formation of the cavities (fig. 6). In crystals showing

the most intense absorption (produced by W. & C. Spicer Ltd., Winchcombe, Gloucestershire from

MgO supplied by Kanto Chemical Co.) the esti- mated concentration of OH" centres (see below)

is ~ lO'^cm"^.

The cavities are not distributed uniformly

throughout the crystal. They appear in a band

whose thickness increases with time but whose

outer edges are at a constant depth below the

surface, irrespective of the annealing time. The

number of lenticular cavities (average diameter

20/xm) within the band in crystals having a

high OH centre concentration is typically ~ 2.5X lO^crn-^. The fact that hydrogen-filled cavities result from annealing in deuterium, carbon monoxide

or hydrogen indicates that the gas in the cavities

stems from an internal source. The presence of

OH" as an impurity in MgO single crystals has

previously been reported by Wertz et al. [8], Kirklin

et al. [9] and by Glass and Searle [10]. These ob- servations, together with the correlation between

the incidence of cavities and the intensity of the

OH" absorption obtained in the present work suggest that the effect of the reducing atmosphere

is to create vacancies for cavity growth and to liberate hydrogen from the hydroxyl ion impurity

of the crystals. Since the presence of hydrogen in the cavities in both slowly and rapidly cooled

crystals indicates that the gas does not precipitate

during cooling, any mechanism for cavity formation

must explain the accumulation of hydrogen during

growth of the cavity. A tentative explanation,

based on the OH"-type centres described by

(a) ( b )

Figure 2. a. Cavity cleaved through its mid-plane.

b. Matched cleavage face after etching to reveal

impurities.

Figure 3. Bright field transmission electron micrograph of an octahedral cavity with {111} facets. Operating reflection

{200).

104
Figure 4. Ultramicroscope photographs showing a possible growth sequence of a lenticular cavity.

Kirklin et al. [9] and by Glass and Searle [10] is as follows. The function of the reducing atmosphere is to remove oxygen atoms from the surface of the MgO,

thereby creating anion vacancies and hberating

Figure 5. Scanning electron micrograph of a lenticular cavity

showing development of lobes, (a, b, c. d)

.

Normal to the plane of the specimen inclined at 35*" to the electron beam.

electrons in the crystal. Wertz et al. [8] showed

that this could occur at 1200 °C even in vacuo (in that work trapping of the electrons by trivalent impurities was considered). The anion vacancies

and electrons then diffuse into the crystal, possibly

in association, until they meet an 0H~ centre where

the following reaction occurs. Mg H0- + 0 + e- Mg 0 + 02- + H. A divacancy containing a hydrogen atom results which may become the nucleus for condensation

of other vacancies to form a cavity. On this basis,

the number of cavities that could be formed would

be related to the concentration of 0H~ ions within

the crystal. The 0H~ centre concentration deter- mined in the present work would provide suf- ficient vacancies for approximately 10"' cavities

per cm^ of average diameter 20/u.m. Furthermore,

the calculated pressure in a cavity assuming that each OH" ion released a hydrogen atom is ~ 500

atm. These figures are in reasonably good agree- ment with the experimental results. The possibility of precipitates of Mg(0H)2 in the

crystals cannot be discounted. Glass and Searle

have associated the absorption band at 3700 cm~'

with the presence of precipitates since this fre- quency is the stretching frequency for 0H~ in Mg(0H)2. (C. Cabannes-Ott. [11].) It may be noted

that the band at 3700 cm~' is prominent in crystals

in which cavities are formed (fig. 6). In the present

290-621 0-68—

8

105
a)

333S 332p

3306|

32931

WAVE NUMBER

Figure 6. Infrared absorption spectra of MgO crystals.

a. For a crystal that does not produce cavities.

b. For a. crystal that does produce cavities; prior to annealing in hydrogen.

c. For the same crystal as b: after annealing in hydrogen.

work one precipitate that gave a selected area

diffraction pattern consistent with crystalline

Mg(0H)2 has so far been found in a thin foil of

MgO. The precipitate was oriented with the c axis

parallel to [111] and the a direction parallel to [110]. This orientation is identical with that deduced

by Garrido [12], Brindley [13] and Ball and Taylor

[14] for the topotactic transformation of Mg(0H)2

to MgO.

3. Cavities Produced by Post-Irradiation Annealing

These have been observed by transmission

electron microscopy (Morgan and Bowen [15]) in the form of cuboids bounded by {100} planes with, a mean length of side of ~ 200 A (fig. 7). They are formed in samples irradiated to doses in excess of

1020 jjyj^ I MeV) by annealing in argon at temperatures above 1500 °C. Cavities are not observed

in similarly annealed samples irradiated to lower

doses or in high-dose samples annealed below

1500 °C. From observations of the size distribution

of the cavities and the fact that, in pulse anneals

at 1800 °C, the cavities formed in periods as short

as 30 seconds, Morgan and Bowen concluded that

the cavities formed by a vacancy condensation

process but that a certain critical vacancy concentration and mobility were required before condensation could begin. In the present work, high precision lattice parameter measurements have been used to gain information about the temperature range over which this condensation occurs and the distribution of the

vacancies prior to the formation of cavities. Measurements were made on a crystal (irradiated to a dose of 8.8 X 10^" nvt) to a precision of ± 2 parts

in 10^ using the Harwell Automatic Precision X-ray

Goniometer (Baker et al. [16]). The results are shown

in figure 8. The initial steep fall in lattice parameter

on annealing between 200 and ~ 1000 °C is due

primarily to the annealing of interstitials, partly

by recombination with vacancies and partly by

clustering to form interstitial dislocation loops on {110} planes (Bowen and Clarke [17]). Approximately equal numbers of anion and cation displacements are to be expected since the fast neutron

cross-sections for Mg^+ and O^" are nearly equal.

Above ~ 1100 °C, the lattice parameter falls below

the value for unirradiated MgO. In this temperature

region, electron microscopy shows that very few

fresh interstitial loops are being nucleated and

106
it is assumed that the majority of interstitials have

akeady migrated to existing loops. It is then rea- sonable to associate the negative change in lattice parameter with lattice dilatations resulting from

the presence of vacancies or vacancy clusters. Condensation of these clusters into relatively large

cavities would result in removal of their dilatations

with a consequent increase in parameter. The temperature at which recovery begins of the negative

change in parameter correlates well with the

electron microscope observation that cavity formation occurs after annealing at temperatures above

1500 °C but it is evident that the process is not complete until the annealing temperature reaches ~ 1750 °C.

The negative change in parameter implies that the dilatations associated with the defects are such

as to produce a net contraction of the lattice. Theoretical (Kemp and Neeley [18]) and experimental (Unruh and Culvahouse [19]) evidence for the F centre in MgO indicates that the nearest neighbour relaxations are outwards and it is possible that the x-ray measurements in this region reflect the presence of vacancy clusters with a net inward relaxation rather than single

vacancies. Martin [20] has deduced, from cold

neutron scattering measurements, the presence

of clusters containing up to ~ 100 vacancies for similar dose samples annealed in the same temperature region.

Helium and neon are produced during neutron

irradiation by transmutation of magnesium and

oxygen ions (Wilks 1966). The total number of

gas atoms generated by a dose of 8.8X10^" nvt, calculated from cross-section data, is ~ 3 X 10'^ cm"^. In contrast to the production of cavities by

OH" ions where the number of hydrogen atoms is comparable to the number of vacancies produced,

the number of rare gas atoms produced by trans- mutation is considerably less than the number of

irradiation-induced vacancies contributing to cavity

formation (~ 3 X 10-' cm"''). Furthermore, gas atoms

are not necessarily produced at the nucleation

sites of cavities. Consequently, the accumulation of gas in the cavities is expected to continue after the

cavities have been formed rather than as an intrinsic part of the cavity formation process. Morgan and

Bowen concluded, from studies of the gas released

by crushing crystals following various periods of anneal at 1800 °C, that despite the presence of

cavities, the majority of gas atoms were initially

trapped at dislocations and that diffusion of ap- preciable amounts of gas into the cavities occurred

only after the annealing period was sufficiently

long to remove the majority of dislocations.

We are grateful to Dr. P. J. Anderson and Mr.

R. D. King for discussions and assistance with

the infrared and mass-spectrometer work and to Mr. T. W. Baker for the lattice parameter measurements. We wish to thank Dr. J. Williams for

his support and interest in this work.

107
NEUTRON — IRRADIATED MgO

4 2I60 ^20 DOSE 8 8x10 n.v.t. ^ I MeV

4 2I50

4 2I40

< 4 2I30

4 2120

4 2110 —

SOO ICOO I500 ANNEALING TEMPERATURE °C

2000

Figure 8. Annealing of lattice parameter changes in irradiated MgO.

4. References

[1] T. Hibi, K. Yada, and Y. Kawamata, J. Phys. Soc. Japan,

18,Suppl. 111,345(1963).

[2] K. Tanaka,~M. Mannami, K. Izumi, and O. Akisue, Japan

J. Appl. Phys. 3, 164(1964).

[3] J. J. Oilman and W. G. Johnston, J. Appl. Phys. 29, 877

(1958).

[4] S. Amelinckx, W. Maenhout-van der Varst and W. Dekeyser, Acta Met. 7,8(1959).

[5] A. D. Whapham, Trans. Am. Nucl. Soc. 8,21 (1965).

[6] H. J. Woollaston and R. S. Wilks, J. Nuclear Materials, 12,305(1964).

[7] D. H. Bowen, Trans. Brit. Ceram. Soc. 62, 771 (1963).

[8] J. E. Wertz, P. Auzins, J. H. E. Griffiths, and J. W. Orton,

Discussion Faraday Soc. 26, 66 (1958).

[9] P. W. Kirklin, P. Auzins and J. E. Wertz, J. Phys. Chem.

Solids. 26, 1067(1965).

[10] A. M. Glass, and T. M. Searle, J. Chem. Phys. 46, 2092

(1967).

[11] C. Cabannes-Ott, Compt. Rend. 242, 355 (1956).

[12] M. J. Garrido, Compt. Rend. Acad. Sci. 203, 94 (1936).

[13] G. W. Brindley, Progress in Ceramic Science, (ed. Burke,

J. E., Pergamon Press) 3, 7 (1963).

[14] M. C. Ball, and H. F. W. Taylor, Min. Mag. 32, 754 (1961).

[15] C. S. Morgan, and D. H. Bowen, Phil. Mag. 16, 165 (1967).

[16] T. W. Baker, J. D. George, B. A. Bellamy, and R. Causer,

Nature, 210, 720(1966).

[17] D. H. Bowen, F. J. P. Clarke, Phil. Mag. 9, 413 (1964).

[18] J. C. Kemp, and V. D. Neeley, Phys. Rev. 132, 215 (1963).

[19] W. P. Unruh, and J. W. Calvahouse, Phys. Rev. 154 , 861

(1967).

[20] D. G. Martin, U.K.A.E.A. Res. Rept. AERE R5521 (1967).

108
Defect Complexes and Microdomains in Nonstoichiometric Oxides. An

Analysis of Wiistite, Fci-sO

G. G. Libowitz

Ledgemont Laboratory, Kennecott Copper Corporation

Lexington, Massachusetts 02173

The deviation from stoichiometry in oxides of the transition elements generally are attributed to an excess of one type of point defect. Consequently, mass transport in nonstoichiometric oxides is usually discussed in terms of mechanisms involving the movement of simple point defects. However,

there is evidence that, at least in some cases, the nonstoichiometry is due to defect complexes and

to clusters of defects (microdomains) randomly distributed in the lattice, rather than to simple point

defects. The evidence for defect complexes or microdomains in nonstoichiometric FeO, UO2, TiO, and rare earth oxides is reviewed. A statistical thermodynamic analysis of defects in wiistite has been carried out and the resulting

expressions relating thermodynamic activity to deviation from stoichiometry were compared with

available oxygen activity data. The results indicate that above 1000 °C the predominant defects are complexes consisting of two Fe vacancies and one Fe*'^ interstitial, rather than simple Fe vacancies. Recent electrical conductivity data are also in better agreement with the defect complex model. Key Words: Defect complexes, microdomains, nonstoichiometry, point defects, statistical thermodynamics, wiistite.

1. Introduction

Nonstoichiometry in oxides usually is attributed

to an excess of one type of point defect. An ap- parent excess in oxygen may be due to interstitial oxygen or metal vacancies; and conversely, an

apparent deficiency in oxygen may be attributed

to oxygen vacancies or metal interstitials. Consequently, mass transport in nonstoichiometric oxides

generally is interpreted in terms of mechanisms

involving the movement of the particular type of

simple point defect which is responsible for the

deviation from stoichiometry. There is evidence,

however, that at least in some cases, the nonstoichiometry is due to defect complexes and to large clusters of point defects, rather than to simple point defects. These should not be confused

with the so-called MagneU compounds [Ij ' which

are separate compounds, closely spaced in composition, whose related crystal structures may be

viewed as made up of different arrangements of

ordered defects. The oxides which we will be

concerned with in this paper vary in composition

over a measurable single phase existence range, and the defect complexes or clusters present are,

to a large extent, randomly distributed in the

crystal lattice. Thus, in order to properly understand mass

transport phenomena in such oxides, the possibility

' Figures in brackets indicate the literature references at the end of this paper.

of the presence of complexes or clusters should be

taken into consideration and the conditions under

which they exist should be evaluated. In this paper,

the evidence for defect complexes and clusters

will be reviewed and the case of wiistite, FeO, will be discussed in detail.

2. Defect Complexes

2.1. Wttstite

It has been shown [2] that wiistite always contains an apparent excess of oxygen; i.e., the stoichiometric

composition, FeO, does not exist. The composition

of closest approach to stoichiometry is FeOi.0.5,

and the deviation from stoichiometry may be as high as FeOi.iH at 1400 °C. x-ray and density studies

[3] have established that there is an excess of Fe

vacancies in the lattice so that the apparent excess

in oxygen is actually a deficiency in iron. A more recent neutron diffraction study of wiistite by Roth [4] revealed that the apparent concentration of vacancies in Fci-sO exceeds the value of

8 as determined by chemical analysis. Also, the

observed diffraction intensity data did not agree

with values calculated on the basis of a model

containing simple iron vacancies. However, the

intensity data were consistent with a model in which defect complexes consisting of two vacancies

and an interstitial Fe ion are present as shown in

109
Figure 1. (a, b and c) — {lll) planes of wustite. (d and e) —

{111) planes of magnetite, Fe^Oi. @ Fe ions on normal sites, # interstitial Fe ions, vacancies, O oxygen anions.

figure 1 (a and b). The defect complexes can cluster

[as shown in (c)] in such a manner as to form

regions which resemble the Fe304 spinel structure

(d and e). Since Fe304 is the oxide which coexists

with wustite at the oxygen-rich side of the phase

diagram, the clusters may act as nuclei for the

precipitation of magnetite when the overall oxygen

content exceeds the saturation hmit of oxygen in

wiistite. Because Roth's studies were on quenched samples of wiistite, there is some question as to whether

these defect complexes actually exist at high temperatures. Recent x-ray studies at elevated temperatures by Koch, reported by Swaroop and

Wagner [5], indicate some sort of clustering does

occur. Further evidence for these defect complexes

will be given later.

2.2. Uranium Dioxide

UO2 exhibits both positive and negative deviations

from stoichiometry. From density and x-ray meas- urements [6], diffusion studies [7] and magnetic

investigations [8], it was concluded that the positive

deviations from stoichiometry (which may extend

up to the composition U02.2n at 1300 °C [9]) were due

to additional oxygen atoms (or ions) situated in the

octahedral interstices of the fluorite structure. However, neutron diffraction studies [lOJ which permit

determination of the positions of the oxygen atoms

showed that, as in the case of wiistite, the situation

in nonstoichiometric uranium dioxide was not that simple. The extra interstitial oxygen atoms do not occupy the octahedral interstices (e.g., 1/2, 1/2,

1/2 positions) but rather are displaced in the (110)

direction to the 1/2, 0.4, 0.4 positions. The two

nearest oxygen atoms are shifted along the (111)

direction from their normal positions about 1 A

towards adjacent interstices. The entire configuration may be considered to be a defect complex

consisting of three interstitial oxygen atoms and

two oxygen vacancies. Here again, the structure of

the next higher oxide phase [11], U4O9, is apparently an ordered arrangement of these complexes, since

the oxygen atoms in U4O9 occur in positions similar

to the ones occupied in the nonstoichiometric

dioxide.

3. Microdomain Theory

Anderson [12] has suggested that defect complexes or clusters such as those described above

may be considered microdomains whose structures

are similar to the next adjacent phase of the com- pound under consideration, and it is the ordering

or aggregation of these randomly distributed microdomains as the composition of the compound

changes which leads to the formation of the new

phase. These microdomains may even exist at the

stoichiometric compositions, such that the lattice of a stoichiometric oxide may contain microdomains

of both the next higher oxide and the next lower

oxide (or metal phase).

On the basis of thermodynamic considerations,

Ariya and Morozova [13] proposed a similar concept

of nonstoichiometric compounds which they re- ferred to as submicroheterogeneity. Ariya and

Popov [14] later suggested a modification of this concept whereby a nonstoichiometric oxide would

consist solely of microdomains resembling the next higher and lower oxides. For example, in TiO,

which has an existence range of TiOo.se to TiOi.20,

the nonstoichiometry has been attributed [15] to oxygen vacancies for oxygen deficient compositions

and titanium vacancies for oxygen excess compositions. Ariya and Popov proposed that the

vacancies cluster into microdomains in the NaCltype TiO lattice as shown in figure 2. In one type

of microdomain, half of the oxygen sites are vacant and the vacancies arrange themselves to resemble

the Ti2 0 structure (hexagonal) although the struc- ture remains cubic. Conversely, the other type of

microdomain has one-third of the Ti sites vacant arranged such that the structure resembles the

orthorhombic Ti2 03 structure, although here again

the actual structure remains cubic. At the stoichio- metric composition, the size of the average Ti203-

type microdomain is about one and one-half times

that of the Ti20-type microdomain. As the overall oxygen content of the sample increases towards

positive deviations from stoichiometry, TiOi + g, the average size of the Ti203-type microdomains

increases at the expense of the Ti20-type microdomains. The converse is true as the overall oxygen

content decreases to TiOi-g compositions. The

110
0 0 T,*3

0 0 0 0 0

0 0 0 0

0 0 0 0 0 0

T,+' 0 0 0

0 0 0 Ti + I 0 D

0 0 0 0 T, + l 0

0 0 0 0 T, + ' 0 T,+l

0 0 0 0 0 0

0 T, +

' D Q 0 T,+3 0 0

0 TI+3 0 Ti+3 0 Ti+3 0 T,+3 0 0

0 0 Ti+5 0 0 T,+' 0 0

0 0 0 0 0

0 T,+5 0 D T,+l Ti+' 0 Ti+' 0

Figure 2. Microdomains in TiO according to Ariya and

Popov [14].

boundaries between the two types of microdomains are continually shifting such that a dynamic

situation exists between microdomains. From the

overall concentration of vacancies [15], Ariya and

Popov have estimated that the average size of a microdomain in TiO ranges from three to nine unit

cells on edge, depending upon the stoichiometry.

Therefore, the microdomains are too smaU to be

seen by normal x-ray techniques which just show

the NaCl-type structure. It may be of interest

that some preliminary electron diffraction studies

[16] of nonstoichiometric TiO (> 900 °C) revealed

diffuse Unes indicative of a disordered complex

superlattice structure, in addition to hues of the

normal NaCl-type structure. Similar type microdomains were also proposed by Ariya for vanadium

monoxide.

The concept of microdomains has been invoked

to explain hysteresis in phase transformation of the

rare earth oxides [17]. This may be illustrated by

referring to figure 3 which shows the partial phase

diagram of the praseodymium-oxygen system and

the results of an isobaric measurement of temperature as a function of oxygen content. The hori- zontal portions of the isobar represent two- phase regions, and the sloping sections repre- sent single-phase regions. It is seen that the ex- perimental isobar exhibits a hysteresis effect across the 2-phase (o-+i) region. The heating

cycle accurately reflects the phase diagram, but on cooUng, the segment representing the two-phase

1.65 1.66 167 1.68 169 1.70 171 172 1.73 1.74 175 176 177 0/Pr RATIO

Figure 3. Partial phase diagram and representative isobar

in the praseodymium-oxygen system {from Hyde, Bevan and

Eyring[17]).

region is no longer horizontal. The cr-phase of

praseodymium oxide is cubic and may be viewed

as a derivative of the fluorite-type structure with

oxygen vacancies ordered such that there are

strings of oxygen octahedrally coordinated Pr^+

ions symmetrically arranged in all four (111) di- rections of the fluorite lattice (the four 3-fold axes).

The i-phase has a similar structure except that it

is shghtly distorted to a rhombohedral structure where the strings of regular oxygen octahedra are now all parallel to only one of the (111) directions

of the pseudo-cube (one 3-fold axes). In passing

from the i-phase to the cr-phase (heating cycle),

these "strings" which are ordered along one (111)

direction become randomly oriented along all four

(111) directions and the transformation is rela- tively sharp as shown. However, in passing from

the disordered cr-phase to the more ordered t-phase

(cooling cycle), there is a chance selection of the one rhombohedral 3-fold axis from the four 3-fold axes of the cr-phase so that different regions of the

crystal have different orientations of the 3-fold

axis. Thus, instead of a sharp formation of the

i-phase, at 975 °C microdomains of the t-phase are formed adjacent with microdomains of the residual

111
cr-phase. As the temperature continues to decrease,

the i-phase microdomains grow at the expense of

(T-phase microdomains and finally at 875 °C, the

transformation to i-phase is complete. A similar

hysteresis effect is observed at higher oxygen contents in the transformation from the rhombohedral

t-phase to a trichnic phase.

Finally it should be emphasized that the microdomain concept is only a hypothesis since the ex- istence of microdomains has only been deduced

from indirect evidence. However, the concept does

provide an appealing explanation of phase trans- formations in nonstoichiometric oxides and is worthy of further investigation. The nature of the

microdomains may be similar to that of the anti- phase structures observed in some intermetallic

compounds such as CuAu [18].

4. Statistical Thermodynamic Analysis

of Wiistite

The remainder of this paper will be concerned

with the defect structure of wiistite. As mentioned

earlier, there is some uncertainty as to whether

simple Fe vacancies or the defect complexes pro- posed by Roth (or even some other types of defects) are responsible for the nonstoichiometry in wiistite. As a further complication, Vallet and Raccah [19]

have proposed, as a result of thermodynamic activity

measurements, that there are three forms of wiistite within the single phase field as shown in figure 4. Despite the phase boundaries shown, the authors

have pointed out that there is a good possibility that some of the phases may exist in a metastable

state within the phase field of others. Additional

evidence for the existence of these phases has

1400

1300

1200

1100

1000

900

800

700

Fey \

1 / *3 /

Fea \ FesOA

1

1

1 1.05 1.10 0/Fe RATIO

1.15 1.20

been obtained from x-ray [20] and dilatometric [21]

studies, but they have not been observed in subsequent activity measurements [22. 23]. No struc- tures have been reported, but it has been sug- gested [24] that the different forms of wiistite are due to various modes of ordering of the defects. Recent interpretations [25, 26] of electrical conductivity data at elevated temperatures 1000 °C)

indicate, on the basis of a mass action law treat- ment, that the defect in wiistite is a doubly ionized

Fe vacancy. However, at the large concentrations

of defects present in wiistite, a simple mass action

law analysis is not valid. More accurate relation- ships for the dependence of electrical conductivity

and thermodynamic activity on oxygen pressure

may be obtained from a statistical thermodynamic

treatment of defects in wiistite.

4.1. Derivation of Equations

The free energy of a wiistite crystal containing an excess of Fe vacancies may be written:

6' = /XpeO + iVnFe^nFe+ OPJh

kT In

(/v-yvnpe)!

(zyVaPe)!(yV-z/VGFe-/VnFe)!j

(1)

where is the total number of Fe sites in the lattice and A^oFe the number of Fe vacancies. /u-Feo is the

free energy of a formula unit of FeO in an ideal

defect-free crystal; ggpe is the free energy of for- mation of an Fe vacancy (exclusive of the contri- bution from configurational entropy) by the addition

of an oxygen atom to the crystal from some standard state (e.g., O2 gas at one atm.); is the average

energy necessary to ionize a positive hole from

an Fe vacancy, and z is the degree of ionization.

If the vacancies are assumed to be doubly ionized,

z becomes equal to two. Since electrical conductivity

in FeO occurs by a hopping mechanism [27], the

holes manifest themselves as Fe+'' ions distributed

among the Fe^- sites in the lattice. The first In term

in eq (1) represents the number of ways of distrib- uting vacancies over Fe sites in the lattice, and the

second In term is the number of ways of distribut- ing positive holes over the remaining Fe atoms. At large positive (excess oxygen) deviations from

stoichiometry, it can be assumed that the concentrations of defects other than Fe vacancies are

negligible. Then, we may write: Figure 4. Partial phase diagram of the Fe-0 system according

to Vallet and Raccah [19]. (2)

112
where Nq and N^^ are the number of oxygen atoms

and iron atoms in the crystal, respectively. If we

consider vacancies to be formed by the addition

of oxygen atoms to the lattice, A^pe remains constant. However, the other quantities in eq (2) change on vacancy formation.

The chemical potential of oxygen in the FeO

crystal, [jlq. is:

dG dG dN,Fe

dNr smce Fe

dNo

dG

1 from eq (2).

(3)

Bearing in mind that the total number of sites does not remain constant on vacancy formation

(i.e., dNIdN^e^l) , we may apply eq (3) to eq (1) using Stirlings approximation to obtain:

Mo ~ f^FeO + ^^DFe + + AT In NiN-SNave)' (4)

Since jjio— fiT In Oq, we may write:

«o ~ Care N{N-3Naver (5)

where Oq is the activity of oxygen relative to the same standard state of oxygen as was used in the

definition of ^pFe- and where

may be incorporated into the FeO lattice. However,

every defect complex in the lattice will block an

additional number of sites or configurations for occupancy by another defect complex. Each va- cancy of the complex restricts twelve possible

vacancy pairs. However, the vacancy pair of the

complex is counted twice so that there are 2(12)— 1 = 23 restricted vacancy pairs. Since each

vacancy pair may be associated with an interstitial Fe ion in two ways, there are 46 restricted or blocked

configurations. Also, the interstitial Fe ion of the

original complex blocks one additional vacancy pair

to give a total of 47 blocked configurations. It can

be shown [28] that the number of ways of distribut- ing Nc complexes over 12A^ configurations with

each complex blocking 47 configurations is:

47-vo mV 47 / ..(f-.,.|: (9)

This expression does not take into account the fact that two or more complexes may exclude the same

site (or configuration). However, this overlapping

effect should only be important at very high defect

concentrations. The free energy of a wiistite crystal

containing the defect complexes may then be

written:

G = NiJLp^o + Ncgc + Nci;,

kT In Wr {N-2Nc)l

(7Vc)!(yV-3/Vc)!j (10)

CoFe = 4 exp [ ( txp^o + MoFe + h)lfiT] . (6)

Since the concentration of Fe vacancies is usually

measured in terms of the amount of excess oxygen

in FeOi+8. we may express eq (5) in terms of 8

utilizing the definition:

8 = /VDFe/(/V- A^QFe) (7)

8-' "'^ =

^°^M-26)^(1 + S)

We now will consider the possibihty that the nonstoichiometry in FeO is due to the defect complexes

proposed by Roth [4]. In the NaCl-type structure,

the coordination number of Fe sites around an Fe

site is 12. Therefore, for each Fe site, there are 12/2

ways of forming a vacancy pair. For each vacancy

pair, there are two ways of placing an Fe ion into an adjoining tetrahedral interstice. Therefore, there

are 12A^ possible ways in which a defect complex

gc is the term corresponding to g^^^ in eq (1) for the formation of a defect complex. is the ionization energy of a positive hole from the defect complex. It is assumed that the interstitial Fe ion is trivalent (i.e., Fe+^), in which case there is only one hole per complex available for ionization. In addition to Wc, the configurational entropy term

includes the term for the number of ways of dis- tributing positive holes over Fe ions on normal

lattice sites. For the case of only defect complexes in the lat- tice, eq (2) can be rewritten as: N^No = Nf~, + Nr

where A^^e represents Fe ions, both on normal sites and in interstices, and eq (3) may be rewritten as:

dG

Mo = ^(11)

113
Application of eqs (11) and (4) to eqs (10) and

(9) yields:

r NUN-2Nr)

- "o [(yv-3/V,) 2 (12/V-47A^e)" "-M12A0

(12)

where

Cc = exp [{fXreo + gc + Ih)lkT] (13)

In terms of deviation from stoichiometry, 8. eq

(12) becomes:

|_(l-2S)^(12-358)"-^«J '^'^^

4.2. Comparison With Data

In order to determine which defect is in better

agreement with experimental data, eqs (8) and (14) were compared with the recent measurements of

Rizzo, Gordon, and Cutler [22] on oxygen activity

in wiistite as a function of oxygen content. If the

standard state of oxygen is taken as oxygen gas at one atm pressure, the oxygen activity, Oq. may be

expressed as the square root of the oxygen pres- sure in atm, PK'. The results for several isotherms 02 above 1000 °C are shown in figures 5 through 9. The

dashed curves were obtained from eq (8) and the

solid curves from eq (14) by computing the values

I

\ 1 \ 1 \ 1 ] 7]—

I

/

/

DEVIATION FROM STOICHIOMETRY, S Figure 5. Comparison of computed activities from eqs. 8 and

14 with oxygen activity data of Rizzo et al. [22] at 1057 °C.

1 r

DEVIATION FROM STOICHIOMETRY, S

Figure 6. Comparison of computed activities from eqs. 8 and

14 with oxygen, activity data of Rizzo et al. \22] at 1075 °C.

of the constants C^pe and Cc which give the best fit with the experimental data. These constants merely

determine the relative heights of the curves. The

shapes or slopes are determined by the particular

function of 8 in each case. It is seen that the data

are in better agreement with the defect complex

model than with the doubly ionized vacancy model

at each temperature. The agreement at 1075 °C

(fig. 6) and for Run #7 at 1150 °C (fig. 8) is par- ticularly striking. The data shown in figure 7 are the result of four separate runs, thus accounting

for the large amount of scatter. However, the ex- perimental points are in consistently better agree- ment with the curve for defect complexes.

The data at 1139 °C (not shown) and 1200 °C

(fig. 9) do not show a clear-cut preference for either

model. In the case of 1200 °C, the highest temperature studied by Rizzo et al., it may be that the

complexes begin to dissociate into single vacancies.

However, recent data of Swaroop and Wagner [5]

show good agreement with the defect complex

model at 1200 and 1250 °C, as shown in figures

10 and 11. In the derivation of eqs (8) and (14), it was implicitly assumed that there was no interaction be- tween defects. If, however, one assumes a pair-wise

Bragg-Williams type of interaction between metal

vacancies (i.e., the number of vacancy pairs is independent of the interaction energy), it can be

114
.05 -06 .07 .08 .09 .10 .11 .12 .13 .14 DEVIATION FROM STOICHIOMETRY , S Figure 7. Comparison ofcomputed activities from eqs. 8 and 14 with oxygen activity data ofRizzo et al. [22] at 1100 °C.

DEVIATION FROM STOICHIOMETRY, S

Figure 8. Comparison of computed activities from eqs. 8 and

14 with oxygen activity data of Rizzo et al. [22] at 1150 °C.

115
10.3

10.4

10.5

°- 106

107

10.8

10.9

1

- -

1 1 1 1 1 1 1 1 1 1 1

1

/-

/

/ /

/ /

/

Voconcles

•/ / //

~ //

- DGf&ct Cornpiex

"^/^ — //

— // // / / / / / / / / / /

/

/ 1 1 1 1 1 1 088 092 0.96 100 .104 .108 .112 .116 DEVIATION FROM STOICHIOMETRY, S

Figure 10. Comparison of computed activities from eqs. 8 and

14 with oxygen activity data of Swaroop and Wagner [5 J at 1200 °C.

9.72

9.80

9.88

rP 9.96-

10-04

10.12

IO20 -

10.28 -,

/ A— Defect Complexes

.092 .096 .100 .101 108 .112 DEVIATION FROM STOICHIOMETRY, 8

Figure 11. Comparison of computed activities from eqs. 8 and

14 with oxygen activity data of Swaroop and Wagner [5] at 1250 °C.

shown [29] that eq (8) will have an additional inter- action term as follows: .

In ao=ln Cnpe + ln

8^ (1-28)2(1 + 8)

. . z£n8(2 + 8)

2kT (1 + 8)2

where is the free energy of interaction between

a pair of Fe vacancies and z is the coordination

number of Fe sites. Thus, the effect, on eq (8), of

neglecting interactions would be an apparent varia- tion of Cofe with 8. Rizzo et al. [22] calculated the

partial molal enthalpies of oxygen, AA/q, in wiistite across the homogeneity range of the compound.

Although there is a great deal of scatter in their

data, there appears to be a slight variation in ^Hq across the existence range. This means that

gfjPe in eq (6), and consequently C^g, do indeed

vary to some extent with 8. From the above rela- tion, we see that the overall effect of interaction will be larger at higher values of 8, and will depend upon

the magnitude of Eq. From activity measurements

at the phase boundaries of wiistite and a simplified

Bragg-Williams treatment, Hoch et al. [30] have

obtained a value of 2.7 kcal/mole for zEqI2. Because

of the assumptions made by Hoch, this is only an approximate value, but it can be used to estimate

the effect of defect interaction on the curves in figures 5 through 11. For the temperatures con- sidered here, log Pq^ would change by only 0.1 from

8 = 0.05 to 8 = 0.14. It can be seen that this would

not lead to a significant difference in the shapes of

the Fe vacancy curves in figures 5 through 11. For the case of defect complexes, estimation of

the effect of interaction between complexes is somewhat more difficult since there will be several

types of defect pairs, each type representing a dif- ferent relative orientation of the two complexes

with respect to one another. Consequently, there

will be a different interaction energy for each type

of defect pair. However, in view of the excellent

agreement of the experimental data with the theoretical curves shown in eqs 5 through 11, it appears

that for defect complexes, the overall effect of defect

interactions is no greater than for the vacancy case. Equations (8) and (14) did not give satisfactory

agreement with experimental data at 1000 °C and

below. This was true for the oxygen activity data of

Ackerman and Sandford [23] and Vallet and

Raccah [31] as well as those of Rizzo et al. Therefore, it appears that the phase diagram proposed

by Vallet and Raccah (fig. 4) is appHcable below

1000 °C, and some other type of defect ordering occurs [24] at lower temperatures.

116
4.3. Conductivity Dependence on Oxygen

Pressures

Recent measurements of the electrical conductivity of wiistite as a function of oxygen pressure

by Geiger at al. [25] and Bransky and Tannhauser

[26] show an approximate pressure dependence of

pi/6 Qvgj- a limited composition range at tempera02 tures of 1000 °C and above, as illustrated in figure 12. The reaction for the formation of a doubly ionized

Fe vacancy may be written:

i O2 ^ + Dpe + 2h^ (or 2 Fe+3) (15)

where Dpe represents an Fe vacancy and /i+ rep- resents a hole or an Fe+'^ ion. The equilibrium constant for this reaction is:

K

[pe][A+]^ Hh-V

Pl/2

O2 Pl/2

' O2 (16)

If the mobility of carriers is constant with composition, the conductivity is proportional to the hole

concentration:

(17)

For the formation of defect complexes, we may

write:

i O, + Fe;^ -> + ( Dpe - Fe,+3 - Dpe ) + ^

(18)

A similar treatment of eq (18) yields a pressure de- pendence of P^^^ for the conductivity. Since the

data in figure 12 indicated a pressure dependence of

P^^^ rather than P^^^ over most of the composition

range studied, the authors concluded, on the basis

of this simple mass action law treatment, that the

predominant defects were doubly ionized Fe va- cancies rather than defect complexes.

However, at the high concentrations of defects

present in this system, a simple mass action law

treatment is not valid. In order to form Fe+'' ions

(or positive holes) on the right hand side of eq (15), Fe+- ions must appear on the left hand side. It is implicitly assumed that the concentration of Fe+^

ions remains constant and that the activity or mole

fraction of Fe+^ is approximately unity in eq (16). However, at the high defect concentrations in this system (5 to 15 percent defects), 10 to 30 percent

Fe^'^ ions are formed and consequently the mole

fraction of Fe^^ ions must vary from 0.9 to 0.7. A

similar objection applies to eq (18). In addition, the

0.6

0.5

0.4

03

0 2

0.1

Bransky a Tannhauser

-1.2 -0.8 -04 0 0.4 08 1.2

2 Log [P(C02)/P(C0)]

1.6 2.0

Figure 12. Log conductivity, cr, and conductance, G, in wiistite as a function of log oxygen partial pressure. The dashed portions of the curves have slopes of 1/6. From Geiger el al. [25] and Bransky and Tannhauser [26).

mass action law treatment does not take into ac- count the possible orientations of the defect complex in the lattice as discussed in section 4.1. A more correct pressure dependence of conductivity may be obtained as follows. Let us assume that the conductivity is proportional to the 1/nth the

power of oxygen pressure, and we wish to deter- mine n. We may write:

^02' °= ^ t^^] ^"1^-

Therefore: In P02 — ^ In (Nn/N) + constant where NnIN represents the defect concentrations

N^rJN or NclN. Hence:

_Nd din Po,

N d(NolN) (19)

Application of eq (19) to eqs (5) and (12) using the

relation ao= reveals that n does not remain

constant at 6 or 4 as imphed by the mass action

law treatment, but varies across the existence range

of wiistite as shown by the calculated values given

in table 1. For the case of doubly ionized vacancies,

the pressure dependence of conductivity varies

from approximately the l/7th to l/9th power. Onesixth power dependence would never be observed

if simple vacancies were present. On the other hand,

the defect complex model leads to l/5th to l/7th

power dependence of conductivity on oxygen pressure. Thus, if defect complexes were present in wiistite, In cr versus In P02 curves should show a

117
Table 1. Variation of oxygen pressure dependence of conductivity of wUstite; a oc PJ/"

Deviation from

stoichiometry 8

n

For doubly

ionized vacancies For defect complexes

0.05 6.7 4.8 0.06 . 6.8 5.0

o:,Q7 7.0 5.2 0.08 7.1 5.4 0.b9 7.3 5.6 0.10 7.5 5.9 0.11 7.7 6.2 0.12 7.9 6.4 0.13 8.1 6.7 0.14 8.3 7.1 0.15 8.6 7.4

sharp slope at low oxygen contents, an approximate

1/6 power dependence of conductivity on oxygen

pressure in the midrange, and a leveling off of the curve at high oxygen contents. This is in accord

with the curves shown in figure 12. Recently, Wagner [32] has obtained some additional conductivity

data and found that there is indeed a continuous

variation in slope for the data in the upper half of

figure 12 rather than a straight line portion as shown. Thus, it appears that conductivity data are

also in much better agreement with the defect complex model than with the model of doubly ionized

simple vacancies.

5. Conclusion

Thermodynamic oxygen activity data and electrical conductivity measurements strongly indicate

that the predominant defects in wiistite above 1000

°C are the defect complexes proposed by Roth [4].

rather than the simple Fe vacancies previously

supposed. Below 1000 °C there appears to be some

other type of ordering in wiistite. Evidence has also been presented that defect complexing and microdomain formation may be quite prevalent in nonstoichiometric oxides. These effects should be con- sidered in the interpretation of mass transport

measurments on nonstoichiometric oxides.

6. References

[1] A. Magneli, Proc. Buhl. Intl. Conf. on Materials, E. Schazt,

ed., Gordon and Breach, New York, 1964, p. 109; A. D. Wadsley, Chap 3, Nonstoichiometric Compounds, L. Mandelcorn, ed.. Academic Press, New York, 1964,

pp. 98-209.

[2] L. S. Darken and R. W. Gurry, J. Am. Chem. Soc. 67, 1398

(1945).

[3] E. R. Jette and F. Foote, J. Chem. Phys. 1, 29 (1933).

[4] W. L. Roth, Acta Cryst. 13, 140 (1960).

[5] B. Swaroop and J. B. Wagner, Jr., Trans. AIME 239,

1215 (1967).

[6] L. Lynds, W. A. Young, J. S. Mohl, and G. G. Libowitz,

Advances in Chem. Series No. 39, 58 (1963).

[7] A. B. Auskern and J. Belle, J. Nucl. Matls. 3, 267 (1961).

[8] A. Arrott and J. E. Goldman, Phys. Rev. 108 , 948 (1957).

[9] L. E. J. Roberts and A. J. Walter, J. Inorg. Chem. 22,

213 (1961).

[10] B. T. M. Willis, Nature 197, 755 (1963).

[11] B. T. M. Wilhs, J. Phys. (Paris) 25 , 431 (1964).

[12] J. S. Anderson, Advances in Chem. Series No. 39, pp. 1-22

(1963); Proc. Chem. Soc. 166 (1964).

[13] S. M. Ariya and M. P. Morozova. J. Genl. Chem. 28, 2647

(1958).

[14] S. M. Ariya and Y. G. Popov, ibid. 32, 2054 (1962).

[15] M. E. Straumanis and H. W. Li, Z. Anorg. Chem. 305,

143 (1960).

[16] J. M. Cowley, Paper presented at Intl. Symp. on High Temperature Chemistry, Argonne Natl. Lab., May 1967.

[17] B. G. Hyde and L. Eyring, Rare Earth Research, vol. 3

(L. Eyring, ed.), Gordon and Breach, New York, 1964,

pp. 623-64; also B. G. Hyde, D. J. M. Bevan and L. Eyring,

Phil. Trans. Roy Soc. (London) 2 59, 583 (1966).

[18] D. W. Pashley, Modern Developments in Electron Microscopy, B. M. Siegel, ed.. Academic Press, 1964, p. 238.

[19] P. VaUet and P. Raccah, Compt. Rend. 2 58, 3679 (1964).

[20] C. Carel, D. Weigel, P. Vallet, ibid. 260, 4325 (1965).

[21] C. Carel and P. Vallet, ibid. 2 58, 3281 (1964).

[22] H. Rizzo, R. S. Gordon and I. B. Cutler, Paper presented

at Symposium on Mass Transport in Oxides, N.B.S., Gaithersburg, Md., Oct. 1967.

[23] R. J. Ackermann and R. W. Sandford, Jr., Argonne Natl. Lab. Report ANL-7250, Sept. 1966.

[24] P. Raccah, private communication.

[25] G. H. Geiger, R. L. Levin, J. B. Wagner, Jr., J. Phys. Chem.

Solids 27, 947 (1966).

[26] I. Bransky and D. S. Tannhauser, Trans. AIME 239, 75

(1967).

[27] D. S. Tannhauser, J. Phys. Chem. Solids 23, 25 (1962).

[28] R. Speiser and J. W. Spretnak, Trans. Am. Soc. Metals 47, 493 (1955).

[29] J. B. Lightstone and G. G. Libowitz, Bull. Am. Phys. Soc. 12, 389 (1967).

[30] M. Hoch, A. S. Iyer and J. Nelkin, J. Phys. Chem. Solids 23, 1463 (1962).

[31] P. Vallet and P. Raccah, Mem. Sci. Rev. Metallurg. 62, 1 (1965).

[32] J. B. Wagner, Jr., private communication.

118
Diffusion in Oxides: Assessment of

Existing Data and Experimental Problems ^

C. E. Birchenall

Professor of Metallurgy, University of Delaware

and

Metallurgist, Lawrence Radiation Laboratory, Livermore, California 94550

Ideal conditions for the direct measurement of diffusion coefficients in oxides are described. The

need for careful measurement, whether the object is to measure intrinsic properties of pure oxides

or the engineering properties of commercial materials, is emphasized. Oxides are classified as impuritytolerant, impurity-sensitive or borderline to judge the likelihood that intrinsic properties have been

measured. A survey of several qxide structure types is given to illustrate the state of the empirical

literature. Key Words: Diffusion, engineering properties, intrinsic properties, nonstoichiometry, thermo- dynamics.

1. Introduction

A clear point, mathematically expressed and

backed by conclusive and consistent data, can be

made quickly. To deal appropriately with the

opposite case will take me a little longer. The

assigned topic poses several problems: (1) The

allocated time could easily be squandered on the

measurements on just one oxide, e.g., BeO. (2) Assessment implies value judgements — that is, do the data satisfy my desires, the desires of the

investigator, or of his employer? Not all of us have

the same concerns. Instead of raising a single, high

ideal, I shall try to anticipate a part of the range

of these concerns. (3) I could exhaust the time and

your patience enumerating and describing experimental methods and the steps that determine the

quality of data. I shall make a few such remarks.

However, most detail is omitted. (4) Finally, if this review stimulates controversy, hopefully even some thought, I shall fulfill another aspect of my

assignment.

The nature of the defects supporting diffusion,

the transport equations, and some ideal model

crystals have been discussed in preceding papers. Assessment of data requires that we inquire what

the actual materials are, intrinsically or extrinsicaUy.

Who needs to know their properties and for what

' Work performed under the auspices of the U.S. Atomic Energy Commission.

purposes? The latter proposition can be dissected

into the scientific approach (what are the properties

of the absolutely pure compound?) and the engineering approach (what are the properties of the stuff

I can buy or make?). Between these legitimate

extremes are many tenable intermediate positions.

The point is that sloppy measurements serve neither

of the ends nor any of the middle adequately.

2. Experimental Methods

2.1. Direct and Indirect Methods

The most direct experiment, that in which only

a single process occurs and in its analysis requires

the fewest unverified assumptions, must be the

standard. For diffusion coefficients, a well-defined

simple initial and final geometry and the direct

verification of composition as a function of distance

and time at temperature are the minimum requirements for a primary standard measurement. The

stated conditions are seldom easy or convenient

to achieve, but in the long run they save time,

controversy and journal pages. Sooner or later someone does it right anyway. In this sense, there

is no standard other than the one described.

Some indirect measurements conform to the

geometrical conditions, but make assumptions

about the spatial distribution of the diffusing

119
substance. The surface activity method and the

'gaseous exchange method, often used for oxygen

diffusion, are of this type. Both of these techniques

are subject to surface transport resistance that may be difficuh to verify and correct for. Vaporization losses can affect the surface activity method.

Exchange with other parts of the vacuum system,

desorption and leakage may affect the exchange

method. Naturally, good technique can minimize

all of these shortcomings, but in the long run the success of the measurement will be judged by how

it agrees with the more direct measurements.

Still less direct measurements result from

inferring diffusivities from processes in which

several mechanisms may contribute simultaneously

to the rate and in which the rate measurement

itself is supposed to distinguish the controlling

mechanism. Oxide growth rates and sintering rates are the most common measurements in this category. Experience indicates that these methods are only reliable under the most ideal circumstances

and careful application.

The invited paper to follow deals with ionic conductivity measurements as a method for studying

transport. The difficulties in correlating diffusivities from conductivity measurement with those from

direct measurement are almost legendary. However,

the combination of these techniques for studying

correlation factors contributes powerfully to a deeper insight into the diffusion mechanism.

2.2. Specimen Preparation

Assume that the direct method of measurement

has been chosen. It consists of several distinguishable steps, the most significant of which is the

preparation of the crystal and its characterization.

The presence and distribution of pores, cracks,

grain boundaries, and impurities should be deter- mined before and after diffusion, or preferably

excluded. For noncubic single crystals, orientation

should be determined and appropriate faces

prepared with minimum distortion. Somewhere

in this sequence the specimen should be equilibrated with a known activity of one of its constituents^ under conditions comparable with those to be used in the diffusion anneal.

Alternatively, the crystal should be enclosed in such a small inert container during the diffusion

anneal that it estabhshes its own equilibrium at- mosphere by dissociation without a significant

change in composition of the surface layer. Packing

the container with chips of the working substance

of identical composition and large specific surface

^ Equilibration with one constituent suffices for a binary system. A ternary system

requires positive control of two activities. For a discussion of this problem with special apphcation to the spinels, see [1].^

^ Figures in brackets indicate the literature references at the end of this paper.

area minimizes superficial changes. The diffusing

substance is estabUshed either by preparing a second crystal of different composition to be welded

to the first or by depositing a thin layer of the

diffusing substance on the surface of the specimen.

Often the deposited substance is a radioactive

tracer to keep the amount small, or to measure

self-diffusion of one of the basic constituents.

Vapor deposition is perhaps the most common

method, but evaporation of a solution containing

the tracer in a compound is also done. Occasionally

electrodeposition may be possible. In any case a uniform deposit to avoid hold-up in clumps is necessary. The necessity to convert the deposit

to the same chemical form as the substrate is not so obvious, but it may be critical for measurements

in all but the most variably composed oxides. Some

attempts should be made to establish the distribution and form of the initial deposit. Microscopic

inspection, autoradiography (for radioactive tracers),

microprobe scanning (for chemically different dif- fusant) are some of the useful methods.

2.3. Diffusion Anneal

The diffusion anneal should be done in a nonreactive system with a uniform and constant hot zone, including good thermal baffling (radiation

shielding). The atmosphere or intimate contact with a second condensed phase should fix the

thermodynamic activity of the specimen throughout

the run. "Vacuum" or "inert flowing gas" are not

satisfactory conditions for establishing a fixed

activity of oxygen over the oxide. Heating and

cooling should not be so slow that composition

changes can occur in these steps, nor so fast that mechanical damage results.

2.4. Sectioning and Analysis

Determination of the spatial distribution of the

diffusing substance after the anneal may be done

in a variety of ways. Slices parallel to the original

surface may be taken by cutting or grinding. Such

cuts must be uniform in thickness. The material

in the slice must be analyzed or the composition

of the exposed surface must be determined. The

amount of material removed in each slice may be

measured or the thickness of the remaining sample

determined. The choices should be taken to optimize

precision. Instead of parallel sections, a single

perpendicular or oblique section may be useful for autoradiographic or microprobe analysis. The

oblique angle gives mechanical magnification of the

penetration distance which may be very desiraoie

when the total penetration by diffusion is small. In most cases it is now better, quicker and cheaper

120
to use the single sectioning method on nonporous

specimens.

Autoradiography and microprobe scanning both

reveal whether diffusion has occurred preferentially

along cracks or grain boundaries, and to what

extent. Autoradiography requires careful calibra- tion of the emulsion if quantitative compositions

are to be obtained. Care must be taken also to correct for absorption of radiation in the specimen

regardless of the detection method. When the

radiation is very weakly absorbed by the specimen

over the penetration distance into the sample, the

detector records the total activity remaining below

it in the sample. When the radiation is very strongly

absorbed, that is, the penetration distance by dif- fusion is much greater than the range of the radiation, the activity measured is nearly proportional

to the surface concentration. Intermediate cases

require more complicated corrections than these

limiting cases. If experimental conditions can be

adjusted to take advantage of the limiting cases, enhanced precision usually justihes extra ettort. Even when parallel sectioning is used, after the

last section is taken in the diffusion zone, it is wise to take an autoradiograph or microprobe scan, whichever is appropriate, in order to look for preferred diffusion paths.

3. Specimen Purity

I have postponed the question of impurities in the specimen because all of the preceding points

affect the utiUty of the data independent of the

purity. The effect of impurities, however, depends

on the purpose of the study and the substance

under study. I shall divide the oxides roughly into two classes — impurity-tolerant and impuritysensitive. Given a choice, concentrate your re- search on impurity-tolerant compounds. These are the substances that can or do deviate more from

stoichiometry than the usual limit below which

the impurity content is readily reduced. A few

other oxides or types of oxides may have mechanisms for tucking away some impurities so that they have little effect on the diffusivities of the

basic ions. The spinels seem to have this capacity

to some extent. It might appear that these impurity-tolerant oxides, because of their variable

composition ranges, require more careful at- mosphere control than the oxides with more sharply

defined composition ranges. Such is not auto- matically the case. The apparent insensitivity of

impurity-sensitive oxides to atmosphere variation

(usually oxygen partial pressure) often has its origin

in the fact that the impurity concentration exceeds

both the intrinsic disorder of the stoichiometric

substance and the composition range of the phase

field of the perfectly pure oxide. Diffusivities ordi- narily vary with oxygen partial pressure to some

power ±l/«, where n generally lies between 2 and 8, frequently between 3 and 6. Thus the

sensitivity to often is not great, and positive

control of P02 is likely to be adequate. Pumping

alone or blowing CO2, H2O or an inert gas alone,

through the system leaves P02 undetermined by

orders of magnitude, which is very undesirable.

For those who must apply impure oxides in a practical device the transport properties in the working substances are the pertinent quantities, not the properties of hypothetical ideal crystals.

Sintered oxide parts are examples where diffusivities on real materials should be useful in understanding and controlling processing. For oxide

growth on shghtly impure metals or on alloys it is important to remember that impurities or alloying elements are fractionated at each interface and

that they are not uniformly distributed through

the growing oxide phases. Application of diffusion

data to such cases requires more than the de- termination of diffusivity in a single composition,

pure or impure. Also oxides for application may

have irreducible variabilities owing to the methods

by which they are produced and handled. If these

factors affect diffusivity the study of the property

must have a statistical sampling basis from the

outset. Tentative examples of oxides that illustrate relative tolerance to impurities are:

Impurity-sensitive: BeO, CaO, MgO, AI2O3

Impurity-tolerant: CoO, FeO, Fe304, cubicstabilized zirconia

Borderline: CdO, ZnO, NiO, aCraOs,

aFeaOa.

The fist includes several oxides that are regarded

as model compounds.

Repeating for emphasis that those who need to know how fast ions move in materials of commercial purity must and should study diffusion in those materials just as carefully as they would

study the purest materials conceivable, I turn to the attempt to find out what the transport properties of pure oxides are — composition variation

within a single field and temperature dependence

particularly. It is largely from these two dependencies that mechanisms of diffusion are deduced.

4. Crystal Structure

The first consideration governing intrinsic trans- port behavior is the crystal structure. I deal mainly

with (1) oxides with the sodium chloride structure

(oxides of the alkaline earths and the iron group

290-621 0-68—9

121
transition metals), (2) the spinel oxides, especially

the ferrites, chromites and aluminates, (3) the

rhombohedral sesquioxides with the a corundum

structure, and (4) cubic stabilized zirconia. You

will recognize that several of my examples are repeat performances, up-dated and amended.

Recall that the structures of many oxides can

be described crudely as nearly close-packed anions

with cations arranged in orderly, or occasionally

disorderly, fashion in the interstices. The description applies even better to many sulfides.

5. Assessing the Data

It is much easier to write directions for a good

diffusion measurement than it is to follow the prescription. I doubt if any set of measurements on

oxides meets all of the criteria I laid down. It is even more difficult to judge the value of a single

set of measurements. You can ask whether they are

internally consistent, if the procedures appear

to have been carefuDy applied, and whether the

magnitudes reported are reasonable. But eventually, assessment of data comes down to the simple

question, can other people get the same results by

similar and by different approaches?

Long ago, when diffusion data on metals were as unreliable as those on oxides are now, I learned

that it is more important to compare actual magnitudes of diffusivities than it is to compare activation

energies. It is more important at this stage to know

that a set of values is approximately right than it is to know that they are internally consistent but

might be off an order of magnitude or so. Also,

when the published data and Arrhenius-type

equation occasionally disagree, the data are usually

right and the equation miscalculated.

5.1. Calcium Stabilized Zirconia

To illustrate, consider calcium-stabilized zirconia, a valuable material because it can be almost ex- clusively an ionic conductor. Figure 1 shows two

sets of data and one Une. The Une is log D vs. IjT

resulting from Carter and Roth's [2] conductivity

measurements converted through the NernstEinstein relation, assuming that the oxide ion

transport number is unity and the correlation

coefficient for oxide vacancies on the fluorite structure is 0.65, which is appropriate for a simple

cubic network [3]. The filled circles are diffusivities determined by mass spectrometric analysis for diffused '^0 in sectioned single and polycrystals

by Simpson and Carter [4]. The open circles were

determined [5] by the time-lag permeabihty method

[6] on sintered tubes of about 95 percent density.

The compositions are not precisely the same in

TEMP - °C

1600 1200 1000 800

10-9 1 1 \ \ I I 1 5.0 6.0 7.0 8.0 9.0 10.0 ioVt

Figure 1. Oxygen diffusivity in calcium-stabilized zirconia. — Calculated [2] from electrical conductivity measurements on Cao.i42Zro,B5«Oi.H58 assuming only oxide ion conduction and a correlation coefficient for anion vacancies

in the fluorite structure of 0.65. 9 "O diifusivities (4) in dense single and polycrystals of Cao.u2Zro.»580i,85» by mass

spectrometric analysis of sections. O Diffusivities from permeability measurements [5] in 95 percent dense tubes of

Cao.08Zro.92O!, 92 by the time-lag procedure [6]. the permeability specimens as in the others. There

are big differences in apparent activation energy.

(Possibly the dissociative gaseous transport across pores at high temperatures contributes to abnormally high diffusivities.) Nevertheless, the whole

picture gives considerable confidence that the

magnitudes of the oxide ion diffusivity are about

right. In making this statement, I do not mean to imply

that the defect structure, the stabilizing mechanism,

or the conditions of stability of calcium-stabihzed

zirconia are understood. For example, the conductivity in the Ume-rich region decreases with in- creasing CaO concentration [7] in spite of the

fact that the vacancy concentration is increasing.

The activation energy for conductivity increases

substantially over this same compositon range. The conductivity decrease may result from the

stabiUty of complexes larger than single vacancies

and lower in mobihty. The structure appears to be

in the early stages of phase separation or ordering

under some conditions that have been investigated.

At 1700 °C the calcium and zirconium ion dif- fusivities are at least six orders of magnitude lower

than the oxide ion diffusivities [8]. Although the

grain boundary diffusivities are about 10^ times

122
greater than the cation bulk diffusivities, the total cation transport should be negligible even in fine- grained material.

5.2. Ion Transport in the Growth of aFe203

Having mentioned grain boundary diffusion, let me emphasize that while the theory of intrinsic diffusion behavior may be advanced most efficiently

by the study of good single crystals, applications

may involve grain boundary transport in an important way. Now that '^O marker studies [9] show that aFe203 grows predominantly by cation diffusion,

it appears that growth rates of Q;Fe203 on Fe304

calculated [10] by the Wagner equation from self- diffusion of iron values are several orders of magnitude too low to agree with the measured growth

rates. Previously this discrepancy did not seem

to be significant.

I do not intend to suggest that grain boundary

diffusion in aFe^Os has been shown to be the growth

mechanism, although that is one possible way to reconcile the discrepancy. However, the aFe203

case history is revealing concerning the state of

knowledge of diffusion in oxides. Lindner [11] first measured iron self-diffusion into a porous, sinteredpowder compact in air. As an afterthought to meas- urements on wiistite (along with a pass at magnetite),

Himmel et al. [10] made one actual measurement

and one measurement that gave only an upper

limit for iron self-diffusion in a natural hematite

single crystal in one atmosphere of oxygen, and the

results were consistent with those of Lindner.

At about the same time, Davies et al. [12] observed

the disappearance of ""Ag markers during the

growth of hematite on magnetite. They attributed

the disappearance of the silver markers to vaporization from the external surface (which was probably

correct), and they concluded that hematite grew

predominantly by oxide ion diffusion inward (which

is probably not correct). Himmel and Holt [9] used

the '**0 activation method [13] to show that oxygen

is practically immobile in Q:Fe203 during its growth.

Hagel [14] measured oxygen self-diffusion by ex- change with single-crystal spheres, and reported

values comparable with those reported for iron

ions in hematite [10, 11]. All of these values are several orders of magnitude too low to explain the

hematite growth rates. Izvekov et al. [15] report iron diffusivities in hematite that are about two orders of magnitude

above the earlier results for similar atmospheric

conditions (fig. 2). These coefficients are still more

than an order of magnitude too low to account for the growth of hematite. Either these results are

still too low in spite of direct disagreement with

earlier measurements, or the iron dififusivity is

TEMP - °C

2000 1400 1200 1000

loV

Figure 2. Self-diffusion in sesquioxides with the alpha corun- dum structure. much greater in hematite in equilibrium with magnetite, or grain boundary transport is mainly

responsible for the transport supporting oxide

growth. However, the growth morphology of hematite is not always that characteristic of simple

volume diffusion control, suggesting that some step

is rate-limiting other than the resistance to diffusion

in some circumstances.

Impurities may have affected strongly the meas- urements on natural hematite. During oxidation

of iron, cationic impurities usually diffuse more

slowly through the lower oxides than does the iron

ion. Fractionation of impurities at the interfaces

also occurs. It appears that Fe203 is normally the

purest of the oxides if the lower oxides intervene

between it and the metallic phase. Whether dif- ferences in purity between growing hematites and

those used for diffusion studies can account for the large discrepancy noted is not known. It is a factor to consider in future work, but it is not yet an explanation for past results.

123
5.3. Diffusivities in aCr^Os and AI2O3

Chromium diffusivities in aCr203 are in substantial agreement in three measurements [16-18]

(fig. 2). They are clearly much greater than oxide

ion diffusivities [19] and are in good agreement

with the growth rate constants as they are related

through Wagner's theory. In Q;Al203 the cation diffusivity [20] is also greater

than the anion diffusivity [21]. Anion diffusivity

shows a strong preference for grain boundary paths.

The growth process of oxide scale on aluminum is so complex that a direct comparison of rates is not appropriate at this time.

5.4. Cation and Anion Self-Diffusion in Oxides

With the NaCl Structure

Figure 3 contains resuhs for cation self-diffusion

in the oxides FeO [10], CoO [22], NiO [23-25],

MgO [26], and CaO [27], each of which crystal- Hzes with the NaCl structure. Oxide ion diffusivities as reported for NiO [28], CoO [29], and MgO [30] are also recorded in the figure. The activation enTEMP - °C

1700 1500 1300 1100 900 800 700

Figure 3. Self-diffusion of cations and anions in some oxides with NaCl structures. ergies Q are obtained by fitting the lines to the

Arrhenius equation

D = D,exp{-QIRT). (1)

Activation energies for the self-diffusion of cations

in FeO and CoO are probably reliable to 1 or 2

kcal/gram-ion. For the other measurements, the

sensitivities of the point defect concentrations to small concentrations of impurities and the sus- ceptibihty of oxygen exchange methods to surface

transport control contributions makes the estimation of uncertainty ranges hazardous. Probably the

activation energy for the self-diffusion of nickel

ions in NiO Hes within the range of values given.

Intrinsic activation energies for the other cases

in the lower left part of the figure may differ from

the quoted values by 20 or 30 percent, which could

result in large deviations in the diffusion coefficients from the reported values. When excess oxygen is added to a stoichiometric

crystal MgO, one filled anion site Oo may be added

to the crystal and one unfilled cation site Vm is created simultaneously. If the temperature is suf- ficiently high, the vacancy may have an excess

negative charge residing on the surrounding anions,

compensated by electron holes that are dissociated

to the extent that they can migrate independently

as current carriers when an electrostatic field gra- dient is imposed.

At high temperatures, it seems to be appropriate

to assume double ionization of at least FeO [31]

and CoO. The other oxides might even follow a dif- ferent pattern because of the difficulty of forming

valence states greater than 2 for Ni, Mg and Ca. Nevertheless, FeO, CoO and NiO are treated here

tentatively as if they formed doubly-ionized cation

vacancies and electron holes. The appropriate re- action equation is

jOA^)^V'i,+ 2i>^ + (yt (2)

where " indicates a virtual doubly-negative charge

associated with the neighboring ions. The virtual charges are not counted in the charge balance,

just as Vm contributes nothing in the mass balance.

The enthalpy change A/// for the formation of one gram ion of cation vacancies also results in the

formation of two gram equivalents of electron holes. For reaction 2 the standard free energy change is

and the mass action constant is given by

/ASO\ / A///\ K= exp ( I exp I - I (3)

When the mass action relation for the reaction is written assuming that the oxide ion concentration

remains substantially unchanged, and that the

concentrations of vacancies and holes are a sat- isfactory approximation for their activities, it is found that the vacancy concentration and the hole

124
concentration are both proportional to the sixth root of the oxygen partial pressure. Also

[F;,] = 4-^/3pi^ exp exp (- ^)

(4) M= 2-p-exp(f)exp(-^)

If the cation diffusivity is proportional to [Vl^^

and the electronic conductivity is proportional to [p+], it is to be expected that measurements of

these properties made over a range of temperatures

at constant composition not too near the stoichio- metric composition will depend only on the activation energy for motion of the corresponding carrier,

A//.W for cation mobility, for hole mobility [32].

If instead the measurements are done at constant

P02 the activation energy in each case must be aug- mented by A////3, because the composition-controlled concentration of carriers is temperature

dependent through eq (4). Measurements that are made at neither constant P02 "or constant composition depend for their utility mainly on the slow

variation implied by p^^^. For other mechanisms of

defect formation, this dependence may be more or

less critical. It is good practice to avoid this question by fixing pressure or composition in addition

to temperature.

Table 1 collects the data for diffusivity of cations

and conductivity by electron holes for FeO, CoO,

and NiO. The data are shown to be consistent by

inserting in the boxes for which no measurements

are available reasonable numbers in parentheses

that are compatible with the measured values. No

stronger claim than consistency can be made be- cause comparisons of reported values and uncertainties quoted by the authors suggest that ranges

of 25 to 30 percent or plus or minus several kilo- calories, whichever is larger, should be associated

with some experimental values. CaO and MgO are omitted from this table because their conductivities

appear to be mixed electronic-ionic in nature, vary- ing with temperature, and very sensitive to small

concentrations of impurities [40]. The increasing

difficulty of oxidizing the stoichiometric crystal, the

reaction implied by eq (2), is seen in the progression of A/// from FeO, which has a negative value,

through a small positive value for CoO, to a larger

positive value for NiO. These numbers measure

the relative stabilities of divalent and trivalent cations in these crystals.

5.5. Diffusion in Some Spinel Oxides

Figures 4, 5, and 6 collect available diffusion data on ferrites (fig. 4, also includes O in C03O4), chroTEMP - °C

1200 1000 800

8 9 ioVt

Figure 4. Self-diffusion of cations in ferrites with the spinel

structure and of oxygen in C03O4.

TEMP - °C

1500 1300 1100 1000

ioVt

Figure 5. Diffusivities of ions in chromites with the spinel

structure and in cobalt titanate.

125
TEMP - °C

1300 nOO 900

O

'

1

' 1

Co in V^C°0.85^'2.10°4 \ AND v \

^^"-^ \

-

- \ \ \^

\ \

-

Cr in NiAl O \ \ \

^

< [53] _ \ \ X^ZnAI.O,

in \^ \

[52] \\\ \

Mg in Mg Al^ 04~^-^<^ Fe in

[51] \\

1 1 1 \ 1 >

^NiAI^O^

\ [53]

ioVt

Figure 6. Cation diffusivities in aluminates with the spinel

structure. mites (fig. 5, also includes Co in two compositions

of cobalt titanate) and aluminates (fig. 6). These

figures are aU on the same scale. They may be compared directly as long as the shift in origin is made

properly. A number of tentative conclusions may

be drawn from these comparisons, but it is important

to note that the experiments are less than ideal in several respects, notably that sintered powder

specimens with residual porosity were often employed, grain boundary contributions were not checked, and the thermodynamic activities of the

components were usually uncontrolled.

Table 1. Energies associated with cation self-diffusion and

p-type conduction in FeO, CoO and NiO. References are in brackets. Estimated values that make the table internally

consistent are in parentheses. All values are in kcaljg mole.

Oxide A//„, + A////3 + A////3

A////3

FeO 30 [10, 22) -10 [10] 10 [19] (-18) -28 [10] -17 [33] CoO 35 [22] (40) (7) 12 [34] 5.1 [22] 46 [23] 1.6 [35] 23 [37] NiO (38) 48 [24] 2.3 [36] 18 [38] 24 [37] 58 [25] 5.5 [37] 21 [39]

In general, diffusion of iron in magnetite is much

faster than it is in the heavily alloyed ferrites. The

iron ion diffusivity is very sensitive to the oxygen

partial pressure [41]. When iron is heavily substituted by ions with less labile valence states, the

cation diffusivities fall to or below the lowest part

of the magnetite diffusion range (that is, for magnetite in equilibrium with wlistite or a low oxygen

partial pressure). The fact that a variety of cations

in a variety of ferrites show roughly comparable

diffusivities (in air) suggests that once iron becomes

predominantly divalent or predominantly trivalent as a result of substituting cations with valence

three or greater or cations with valence two or smaller, respectively, further composition change

is not large in its effects [42].

Cobalt titanate [50] has cobalt diffusivities comparable to the cation diffusivities in alloyed ferrites. They show a strong dependence on the ratio of cobalt to titanium but very little dependence on the

oxygen partial pressure [43]. Perhaps the strong

effect of oxygen partial pressure on cation mobility

in magnetite is through its effect on the ratio of Fe-+ to Fe^+ and not directly on the cation to anion

ratio. Clearly the general mobihty of cations in ferrites

is greater than that in chromites which in turn is sHghtly greater than that in aluminates. This trend

agrees with the relative growth rates of these products on ferrous alloys at high temperatures; that is, alloying with chromium reduces the oxidation rate of iron-base alloys that form predominantly spinels,

while alloying with aluminum reduces the rate even more. The activation energies for individual

materials are probably not very significant. If a rough average slope is taken through each set of data, a characteristic activation energy is found for

ferrites to be about 67 kcal, for chromites to be

about 88 kcal, and for aluminates to be about 92

kcal/ gram-mole of cations.

Diffusion coefficients of the two cations in the same spinel may be quite different. Although the

data are not totally consistent, it appears that D^o is greater than 0*^ in cobalt chromite and that the difference increases with increasing ratio of

Co to Cr. D*. appears to be lower than D*^ in NiCr204, and D*^ appears to decrease as the ratio of Cr to Ni increases. Oxygen diffusivity in NiCr204

is lower than either cation diffusivity, but it may

not be as low as the diffusivity of iron in this substance. These wide mobihty differences explain

in part the nonuniform distribution of the metallic

elements in scales on alloys. They also suggest

the possibihty that scales might exist in which cation

dominance of growth gives way to anion dominance

within the scale layer because of composition

changes. However, I know no real example at this time. In nickel aluminate, the chromium diffusivity

exceeds the iron diffusivity which in turn exceeds

the nickel diffusivity.

126
6. Conclusion

I have not cited at all several oxide systems that

illustrate good agreement among investigators,

as in CuoO, or poorer agreement, as in ZnO. I have

overlooked the pseudo-oxides, such as AgS, which

have such high intrinsic disorder in the cation sites that cation mobilities have httle concentration

dependence and very low temperature dependences. Nor have I discussed another pseudo-oxide,

FeS, that illustrates the demobilizing effect of a vacancy superlattice in the cation sites at low

temperatures at about composition FeySg.

Nevertheless, I think I have represented fairly

the state of this art by the examples chosen. There

are some measurements that are rehable because

of the care with which they were done or the confirmation they have subsequently received or both.

Other measurements must be regarded as tentative

until independent confirmation of a direct or cor- relative sort develops. Correlations of the type I have offered will become increasingly useful and

suggestive as the amount and range of the data in- crease. Some of the values I presented wiU be

omitted from the later compilations because they

failed to stand the test of repetition, correlation or consistency with the better theories that we hope

will be available in the near future.

7. References

[1] H. Schmalzried and C. Wagner, Z. Physik Chem. N.F. 31,

198 (1962).

[2] R. E. Carter and W. L. Roth, unpublished research quoted

in reference [4]. [3] C. Compaan and Y. Haven, Trans. Faraday Soc. 52 , 786

(1956).

[4] L. A. Simpson and R. E. Carter, J. Am. Ceram. Soc. 49,

139 (1966).

[5] A. W. Smith, F. W. Mezaros and C. D. Amata, J. Am. Ceram.

Soc. 49 , 240 (1966).

[6] R. M. Barrer, Diffusion in and Through Solids, Cambridge

University Press, 1941, pp. 18-19.

[7] J. M. Dixon, L. D. LaGrange, U. Merten, C. F. Miller and

J. T. Porter II, J. Electrochem. Soc. 110, 276 (1963); T. Y. Tien and E. C. Subbarao, J. Chem. Phys. 39, 1041

(1963).

[8] W. H. Rhodes and R. E. Carter, J. Am. Ceram. Soc. 49,

244 (1966).

[9] L. E. Himmel and J. B. Holt, Chicago Meeting, AIME,

October, 1966.

[10] L. E. Himmel, R. F. Mehl and C. E. Birchenall, Trans. Met.

Soc. AIME 197 , 827 (1953).

[11] R. Lindner, Arkiv Kemi 4, 381 (1952).

[12] M. Davies, M. T. Simnad and C. E. Birchenall, Trans. AIME,

J. Metals 3, 889 (1951); 5, 1250 (1953).

[13] R. H. Condit and J. B. Holt, J. Electrochem. Soc. Ill, 1192

(1964). [14] W. C. Hagel, Trans. Met. Soc. AIME 2 36, 179 (1966).

[15] V. I. Izvekov, N. S. Gobunov and A. A. Babad-Zakhrapin,

Fiz. Met. i Metalloved. 14, 195 (1962).

[16] I. M. Fedorchenko and Y. B. Ermolovich, Ukrain. Chim.

Zhur. 26, 429 (1960).

[17

[18

[19 [20

[21 [22

[23

[24;

[25

[26

[27

[28

[29

[30

[31

[32 [33

[34

[35

[36

[37 [38

[39:

[40

[41 [42

[43

[44:

[45

[46

[47: [48

[49

[50

[51

[52

[53

[54:

[55

W. C. Hagel and A. U. Seybolt, J. Electrochem. Soc. 108,

1146 (1961). L. C. Walters and R. E. Grace, J. Appl. Phys. 36, 2331

(1965). W. C. Hagel, J. Am. Ceram. Soc. 48, 70 (1965). A. E. Paladino and W. D. Kingery, J. Chem. Phys. 37, 957

(1962). Y. Oishi and W. D. Kingery, J. Chem. Phys. 33, 480 (1960). R. E. Carter and F. D. Richardson, Trans. Met. Soc. AIME

200, 1244 (1954).

J. S. Choi and W. J. Moore, J. Phys. Chem. 66, 1308 (1962). M. Klotsman, A. N. Timofeev and I. Sh. Trakhtenberg,

Fiz. Met. i Metalloved. 14, 91 (1962). R. Lindner and A. Akerstrom, Disc. Faraday Soc. 23, 133

(1957). R. Lindner and G. Parfitt, T. Chem. Phys. 26, 182 (1957). R. Lindner, St. Astrumdal and A. Akerstrom, Acta Chem.

Scand. 6, 468 (1952). M. O'Keefe and W. J. Moore, J. Phys. Chem. 65, 1438

(1961). The line is calculated from the corrected equation

given on p. 2277. W. K. Chen and R. A. Jackson, Am. Ceram. Soc. BuU. 46,

357 (1967). Y. Oishi and W. D. Kingery, J. Chem. Phys. 33, 905 (1960) and 34, 688 (1961).

I. Bransky and D. S. Tannhauser, Trans. Met. Soc. AIME

2 39, 75 (1967). The AH^ value in table 1 is calculated from conductivity values for Feo.ssO. C. E. Birchenall, Met. Rev. 3, 235 (1959). K. Hauffe and H. Pfeiffer, Z. Elektrochem. 56, 390 (1952), and Z. Metallkunde 44, 27 (1953). C. Wagner and E. Koch, Z. Physik Chem. 32B, 439 (1936). R. Uno, J. Phys. Soc. Japan 22, 1502 (1967). His measure- ments are consistent with singly ionized vacancies instead

of the doubly ionized vacancies assumed. That is, conductivity is proportional to p^^-^.

F. J. Morin. Phys. Rev. 93, 1199 (1954).

S. P. Mitoff, J. Chem. Phys. 35 , 882 (1961). H. H. von Baumbach and C. Wagner, Z. Physik Chem.

24B, 59 (1934). M. G. Harvk^ood, N. Herzfeld and S. L. Martin, Trans. Fara- day Soc. 46, 650 (1950). S. P. Mitoff, J. Chem. Phys. 41, 2561 (1964). H. Schmalzried. Z. Physik Chem. N. F. 31, 184 (1962). In addition to the data given, some unpublished results in my laboratory support this conclusion. V. I. Izvekov, Inzh. fiz. Zh. Akad. Nauk. Belorus. SSR 1, No. 11, 64 (1958). R. H. Condit, M. J. Brabers and C. E. Birchenall, Trans. Met. Soc. AIME 218, 768 (1960). R. H. Condit, M. J. Brabers and C. E. Birchenall, unpub- hshed data. R. Lindner, D. Campbell and A. Akerstrom, Acta. Chem.

Scand. 6, 457 (1952). R. Lindner„^Arkiv Kemi 4, 381 (1952). Barbara A. Thompson, Ph. D. Thesis, Renssalaer Poly- technic Institute, 1962. R. Sun, J. Chem. Phys. 28, 290 (1958). A. Morkel and H. Schmalzried, Z. physik Chem. N. F. 32,

76 (1962). R. Lindner and A. Akerstrom, Z. physik Chem. N. F. 6,

162 (1956). R. Lindner and A. Akerstrom, Z. physik Chem. N. F. 18,

303 (1958).

I. N. Belokurova and D. V. Ignatov, Soviet J. Atom. Energy

4, 399, transl. p. 301 (1958). W. D. JGngery, D. C. Hill and R. P. Nelson, J. Am. Ceram.

Soc. 43, 473 (1960). N. S. Gorbunov and V. I. Izvekov, Uspekhi. Fiz. Nauk 72,

273 (1960). 127

The Determination of Thermodynamic Properties in Single Phase Wiistite

by Coulometric Titration in a High Temperature Galvanic Cell

H. F. Rizzo\ R. S. Gordon, and I. B. Cutler

Department of Ceramic Engineering

University of Utah

Salt Lake City, Utah 841 12

The thermodynamic properties of wiistite were determined as a function of composition between

1050 and 1200 °C by coulometric titration in a high temperature galvanic cell using a stabilized zirconia

electrolyte. Isothermal plots of the oxygen pressure versus composition (0/Fe) were sigmoidal in shape

with a pronounced negative deviation from Raoult's Law below the inflection point (FeO1.10-FeO1.12) and a positive deviation above. The partial molal enthalpy of oxygen, A//o, varied between —64.5

kcal/mole (FeOi.os) and —74.5 kcal/mole (FeOi.14). A statistical mechanical model, developed by

Libowitz and based on Fe + ^ interstitials, each associated with two Fe vacancies, predicted correctly the P02 versus O/Fe dependence at temperatures between 1057 to 1200 °C. The usefulness of a sohd

electrolyte galvanic cell in determining phase boundaries and thermodynamic functions in the iron- oxygen system was demonstrated. Key Words: Composition-versus-pressure, defect complexes, galvanic cell, wiistite.

Introduction

The high temperature galvanic cell is especially

suited for studying the thermodynamic properties

of defect oxide phases. By varying the composition

electrochemically (i.e., coulometric titration) the

activity of oxygen can be obtained as a continuous

function of composition over the entire phase field of a nonstoichiometric phase. Coulometric titra- tion provides the experimental flexibility and control necessary to delineate the composition and

phase boundaries in a number of oxide systems.

With respect to the study of defect phases, two

basic techniques have been employed.

1. Measurements have been conducted on nu- merous samples of a defect phase, each pre- pared at a specific composition [1-5]."

2. Measurements have been made on a single

sample, in which the composition is varied

electrochemically by a coulometric titration technique [1. 6-8].

The first technique has been used to study regions

of homogeneity in the defect phases Fcj-O [4], WO2-X [5] , Ti02_^ [3] , UO2-U3O8 [1] , and Nh-.O,-^

[2]. The second technique has been used by Kiukkola and Wagner [7] to determine single and two-

' Presently Staff Scientist (Ll. Col., USAF), Air Force Weapons Laboratory, Kirtland Air_Force_Base, New Mexico 87117. ^Figures in braclcets indicate the literature references at the end of this paper.

phase regions in the Ag-Te system. It has also been

used in the UO2+X [6], UO2-U3OH [1], TiOo-x [8],

Nb02±x and Tio.5Nbo..502^x [8] systems.

The objective of this study was to determine the

feasibility of coulometric titration in determining

phase boundaries and oxygen activities in oxide

systems. The iron-oxygen system was selected since

there is considerable information, principally the

work of Darken and Gurry [9] , on the phase

boundaries and thermodynamic properties of the

wiistite phase as a function of temperature and

composition.

Experimental

Two types of experimental cells were employed

in the studies: (1) a closed-end tube cell and (2) a three-electrode cell. The closed-end tube cell consists of a dense stabilized zirconia (.85 ZrOa + -15 CaO) tube (~ 1" OD by ~ 10" long) closed at one

end. (See fig. 1.) The inside of the tube is open to the

atmosphere (Salt Lake City, Utah) which is used as the standard reference oxygen potential (^03(31^) = 0.18 atm). The other side is sealed in a chamber

in which the atmosphere is purified argon. A detailed description of the closed-end tube cell and the

fabrication of the electrolyte can be found else- where [10]. With the foregoing type of cell arrange- ment, three different compositions were studied:

129
Pt, Fe-FcxO

Pt, Fe^O-Fe304

Pt, FeOy

0.85 Zr02 +

0.15 CaO

0.85 Zr02 +

0.15 CaO

0.85 Zr02 +

0.15 CaO

O2 (air, F" =0.18 atm), Pt

O2 (air, P" =0.18 atm). Pt

O2 (air, P" =0.18 atm), Pt

Spring Loaded

Mullite Tube

III

The anode compositions FeOy and Fe^O-Fe304 were obtained by titration of oxygen into Fe-FejO

sample electrodes.

The cell emf (E) for cell I, II, and III is related

to the equilibrium partial pressures in atmospheres

of diatomic oxygen at the anode (^q^) and cathode

{PqJ by equation (1), in which R is the gas constant, is the Faraday constant and

02

(1)

T is the absolute temperature. Using cells I and II the oxygen activity {P'o^) was measured for the two

phase mixtures, Fe-Fcj-O and Fe^O-Fe304 (II). In

addition, the oxygen activity was measured (using

cell III) for any arbitrary composition (y) in the

wiistite phase (FeOy) by coulometric titration. The three-electrode cell (fig. 2) consists of three

electrodes in a stabihzed zirconia cylinder (1 in.- diameter, in. -height). In the coulometric titra- tion studies, one of the bottom Fe-FcjO electrodes was used to titrate oxygen in or out of the top elec- trode, while the other bottom electrode was used

to record the emf. Two compositions were studied

in this cell:

Pt, Fe-Fe^O

Pt, Fe-FexO

0.85 ZrOa +

0.15 CaO

0.85 ZrOz +

0.15 CaO

Fe,0-Fe304, Pt;

FeO„

z<\ly<x

IV

V

Isothermal titration experiments were conducted

in both directions with cell V across the entire

wiistite phase between 800 and 1200 °C. In order to calculate the activity of oxygen at the cathode

(FezO-Fe304), the activity at the anode must be

known. The oxygen potentials for the Fe-Fe^O

equilibrium are reasonably well estabhshed [9,

11-15] and were also measured in this study [16]

with the closed-end tube (cell I). In the titration studies the cell voltages and cur- rents were measured on a cyclic basis. Using an

Gas Outlet

Mullite Protection Tube

Pt Electrode & TC

Alumina

Disc

Alumina Disc

Pt Electrode 8TC

Fe + Fe^O Sample

Electrode

Figure 1. Schematic of galvanic cell with zirconia electrolyte closed-end tube cell.

Gas Outlet

Mullite Protection Tube

Alumina

^ The compositions in the phase field of wiistite are signified FeOy where y is the O/Fe ratio.

Alumina Disc

Pt Electrode & TC

Wustite Sample

Electrode

Zirconia

Electroylte

Alumina Disc

Figure 2. Schematic of galvanic cell with three-sample

electrode, cylindrical shaped zirconia electrolyte.

electromechanical timer the cell voltage was re- corded automatically after each titration period.

The normal procedures in the titration experiments were as follows: (1) Start titration from the

iron-rich boundary of wiistite; (2) vary the titration current for a fixed titration time so that the time

after titration required for the cell voltage to stabihze at a constant emf (equihbrium) would be

within the time allowed; (3) reverse titration direction and compare the ratio of change in emf per

130
coulomb titrated in both directions; (4) periodically

interrupt titration cycles and measure cell voltage

for extended periods (1-24 hours). Analysis of data

from items (1) through (4) was the basis for deter- mining if the titration experiment was a reversible

process and if equilibrium cell voltages were being

measured. Measurement of cell vokages during the

titration period provided information on the degree

of polarization. Titrations in both directions were

performed between the wiistite boundaries. Minimal hystersis and equality of the number of cou- lombs titrated in both directions were indications

of rehable experiments.

One of the principal experimental difficulties was the reduction or oxidation of the single-phase

wiistite sample by the cell atmosphere. The ap- parent rate-controUing mechanism was the transfer

of oxygen from the surface of the sample to the gas phase. Reduction and oxidation rates could be

changed- by varying the gas flow rates and gettering

conditions. Verification of the gas atmosphere as a controUing factor was determined by monitoring

the cell atmosphere surrounding the wiistite elec- trode by the use of a platinum electrode on the

zirconia tube next to the wiistite electrode. It was

noted that during titration of oxygen into wiistite the

oxygen pressure increased in the cell atmosphere

and decreased during titration of oxygen out of the

wiistite electrode. The reduction of the wiistite electrode was also noted in the three-electrode cell and was most probably due to the presence of the

two Fe-FexO sample electrodes.

Corrections for reduction or oxidation of the

wiistite electrode were made on a cyclic basis. Two methods were used. One was to measure the

rate of change of cell voltage for extended times

(45 min. to 24 hrs.) for a known change in the number of coulombs (i.e., composition). From this measurement, the change in the number of cou- lombs per unit time for a given voltage range was

calculated and used in each titration cycle to estimate the change in the number of coulombs due

to oxidation or reduction taking place during the

emf measuring period. The second method was to compare the change in the cell emf during a titra- tion step with the net change in cell emf between

adjacent cycles. If these voltages were different, an effective number of coulombs to account for the

difference was either added to or subtracted from

the number passed in the titration period. The

details describing these corrections are found in the thesis of the principal author [10]. The magnitude of these corrections in estabUshing the oxygenpartial pressure-composition isotherms for the

single-phase region varied between three and

twenty percent.

The electrode samples used in this study varied

from pure Fe (less than 25 ppm) to samples containing 77.0 weight percent Fe. Compositions containing 85 and 82 weight percent iron were pre- pared by dry mixing and pressing mixtures of

reagent grade Fe^Os (99.0 percent minimum) and

iron powder (GAF HP carbonyl 99.6 to 99.9 percent)

in V2 inch steel dies and then heating in a tungsten vacuum furnace (2 X 10^^ mm Hg) between 950

to 1000 °C for 19 hours. A sample containing 77 wt

percent iron was prepared by oxidation of iron

(99.96 percent) in a H2O/H2 atmosphere at 900 °C. This powdered sample was used in the three-elec- trode cell for the variable composition wiistite. The weights, dimensions, and total iron deter- mined by chemical analysis for the electrode

samples used in the titration experiments are pre- sented in table 1.

Results

A. Fe-FcjO, FeeO-Fe304 Phase Boundaries

In the progress of the titration experiments on

cells I, II, and IV, the oxygen potentials were deter- mined for both the Fe-FcjO and FejO-Fe304 (i.e..

Table 1. Summary of electrode samples used in titration experiments

A B C D E

Composition of electrode (wt % Fe) 85.5 82.0 100 ' 77.0 85.5 Weight of electrode (g) 0.221 0.2222 0.3385 0.0922 0.2336 0.265 Sample dimensions (cm) 1.5 length 1.1 diam. 0.6 length powder powder

0.26 width 0.070 thick 0.3 width compact compact

0.10 thick 0.065 thick 1.5 diam. 0.6 diam.

0.023 thick Surface area (cm^) 0.39 0.95 0.18 1.77 0.28

A, B, C Electrode sample — closed-end tube D Electrode sample — three-electrode cell E Reference electrode — three-electrode cell

131
composition 0/Fe~1.33)^ two phase equilibria.

The experimental details and results are reported

extensively elsewhere [10, 16]. For the purpose of

this paper the standard free energy of formation

(calories) of Fcj-O and log Pq^ for the Fe-Fe^O equilibrium {x Fe+1/2 02 = Fej-0) are given by equations (2) and (3) and are valid between 540 and

1200 °C [10, 16]:

AF2 = - 62,452 ± 46 + 15. 127 (± 0.05) r(°K): (2)

-27,295 ±85

logPo^^ f

h 6.6115 ±0.0218. (3)

Equation (4) is an expression for log P02 for the

Fe^O-Fe304 equilibrium between 900 and 1200

°C [10, 16]:

log Po, = ~ ~

+ 13.354 ± 0.064. (4)

B. Titration in Single Phase Wiistite

Two principal techniques were employed: (1) At

a constant temperature oxygen was titrated into the wiistite phase for a prescribed time, followed

by measurements of the cell emf for a fixed time

sufficient for equilibrium to occur. Repeated ti- tration-emf measurement cycles resulted in the

estabhshment of an oxygen pressure-composition

isotherm. (2) The same basic procedure was fol- lowed, but after each titration period an emf-temperature profile was made for each new composition, giving the oxygen potential as a function of

both temperature and composition. The latter method proved more difficult because of the

tendency for the single-phase wiistite to become

oxidized or reduced by the surrounding cell atmosphere.

The reaction for cell III with the oxygen pressure

corrected to one atmosphere is given by equation (5):

1/5 FeOy+1/2 O2 (1 atm)=l/5 FeOy+s y> > >5.

(5)

The equilibrium oxygen pressure for any wiistite phase FeOy is related to by equation (6):

* Evidence, which is reported elsewhere [17], was found in the titration experi- ments of this study that several defect phases exist between the presently reported oxygen-rich boundary of wiistite and mafinetile [9]. Therefore, reference to the Fe^O- FejOi equilibrium in this paper corresponds to the most oxygen-rich wiislite-type phase in equilibrium with magnetite or a composition slightly less oxygen rich than FeO. M (i.e., FejO,).

The cell reaction for the three-electrode zirconia

cell (V) is given by equation (7):

xFe + ^^^Fey'0 = FexO + ^Fey'+„0. (7)

By subtracting the reaction (xFe+1/2 O2 (1 atm) = FcjO) from equation (7) the activity of oxygen in wiistite can be determined, assuming that y' > u. The equilibrium pressure of oxygen computed from

Ey and AF/ of Fcj-O is given by equation (8):

- log Po. =

2 303 RT^~ ^"^^^ '^^^ - 15. 127r]

.

(8)

Oxygen pressure-composition isotherms were

determined in coulometric titration experiments

at 800, 894, 950, 1000, 1057, 1075, 1100, 1139,

1150, and 1200 °C. Selected isotherms are presented

in figures 3-6 for the temperatures 1075, 1100,

1139, and 1200 °C, respectively. One interesting

feature of all the Pq^ vs. 0/Fe plots is their sigmoidal

shape. The correlation of the inflection points in the curves with the p-n transition (represented

by the arrows on figures 3-6) in wiistite will be disToe ' Toe ' Tio n2 m4

O/Fe

Figure 3. 1075 °C isotherms of P^^ versus composition, O/Fe.

In Po,= '^^f,

"' + ln (0.18). (6)

132
1200-

1 1000

Q 800

600

40C200

^ocUJqU

T= IIOO'C ° V ^,v^M!

= Cell m. Sample CTitratkyi Out ocoOQ3°o° V = Cell*, Sample B, Titration In ' o = Cell V, Sample D, Titration In

o^ooooo pO°o

9? °

a

J I L_ _i I ; i_ 106 108 110 I 12 114 116 1 18 I 20 I 22

0/Fe

I 26

Figure 4. 1100 °C isotherms ofP„ versus composition, O/Fe.

700-

100

o Cell V , Sample 0, Titration In, II39°C a Vol let a Raccoh II40°C

_1 I I L

I 07 III I 15 I 19 1 23 I 27 0/Fe Figure 5. 1139 °C isotherms of P^^ versus composition, OjFe.

cussed later. Plots of log P02 versus composition were employed in the Gibbs-Duhem graphical

integrations. The raw data and the experimental

conditions used to construct the isotherms are tabulated in the thesis of the principal author [lOJ. The compositions were determined from a knowledge of the weight of the iron in a particular sample

and the assumptions that the composition of the

iron-rich boundary was FeOi.osn and was constant

Cell V, SompleD

Titrotlon In

Figure 6. 7200 °C isotherms of Pq^ versus composition, OjFe.

in the temperature range of measurements. The

validity of these assumptions is discussed else- where [10, 16]. In general the oxygen-wiistite composition iso- therms presented in figures 3 through 6 indicated

higher partial pressures than those previously

reported for the single-phase region, except in a few cases. This discrepancy is difficult to explain

unless the cell voltage measurements did not cor- respond to equilibrium conditions. Nonequilibrium

voltages could result from (1) side reactions be- tween the electrolyte and the wiistite, (2) a reaction

between the platinum electrode and wiistite, (3) insufficient time for diffusion to occur after titra- tion, and (4) polarization of the cell. The first two can be dismissed when one considers the reproducibility of measurements in the

133
two-phase regions Fe-FcxO and Fe20-Fe304 combined with the excellent agreement of the partial

pressures of oxygen with other investigations [10]. Polarization of the cell could occur at low temperatures and high titration currents. For example,

in establishing a partial pressure composition iso- therm at 800 °C, titration currents greater than

0.4 milliamps required times longer than the allowed

35 minutes for stabilization of the voltage. In the course of a titration experiment the voltage drops

across all components of the cell were measured

and titration currents were reduced whenever

there was an indication of polarization at one of

the sample electrodes. The fundamental experimental criterion used to establish that equilibrium

measurements were being made was that the cell voltage was time-invariant. Hysteresis effects due

to direction of titration were small and the re- producibility of oxygen pressure composition iso- therms was good. For example, the 1100 °C

isotherms (fig. 4) were obtained from four separate

experiments which employed different cells,

samples, and titration currents (0.4 to 2.1 milli- amps), titrations in both directions, different times

to measure cell emf after titration (9-35 min),

varying sample thickness and starting compositions.

The composition where the first two-phase region

occurred (constant oxygen activity) compared to within ± 0.005 0/Fe of the reported oxygen-rich

boundary of wiistite (FeOi.154). This agreement

confirms the validity of determining the actual

phase boundary composition by coulometric

titration. The series of plateaus for compositions

greater than Fed.154 correspond to the multiphase

region previously mentioned [17]. The temperature

of 1100 °C was certainly high enough to dispel the

question of diffusion and, although there was

scatter between experiments there is little doubt

that the partial pressure of oxygen was higher than

reported values. It is noted that all reported data were obtained with CO-CO2 or H2O-H2 atmospheres,

so the effect of the presence of carbon or hydrogen

on the activity of oxygen in wiistite could perhaps

provide the explanation for the lower partial pres- sures measured by these methods. The data re- ported by Barbi [4] on wiistite, which were

determined by emf measurements, also showed

higher oxygen pressures than those reported by

others. The difference in partial pressure of oxygen be- tween this study and other investigations is even more pronounced for isotherms obtained between

800 and 1000 °C. This observation is again difficult to explain since in this temperature range there was very little oxidation or reduction of the sample

and thus, if diffusion was not a rate-limiting step

in achieving equilibrium, the measurements should

be valid. However, it was observed many times

throughout this investigation that the emf-temperature coefficient would undergo a change in slope be- tween 800 and 950 °C for all compositions more

oxygen-rich than FeOi.oe- It is also noted that Kleman [12] has proposed the existence of more

than one wiistite defect phase in this temperature

and composition range. Therefore, to avoid the

question of nonequilibrium isotherms and the possibility of other defect phases being present, only

the isotherms between 1057 and 1200 °C were employed to calculate the thermodynamic properties of wiistite. The data at the lower temperatures can be found elsewhere [10]. It will also be shown

that use of a defect model for wiistite proposed by

Libowitz [30] supports the use and validity of the

higher temperature isotherms.

C. Thermodynamic Treatment of Wiistite

From the relationship between the pressure of

oxygen and the composition of wiistite at various

temperatures the partial molal thermodynamic

functions can be calculated. The partial molal

quantities are particularly useful for describing the

variation of thermodynamic quantities with the

deviation from stoichiometry. The choice of components is somewhat arbitrary. For example, the

following reactions describe the formation of

wiistite:

Fe+(y/2)02 = FeOy5; (9)

FeOi.o5 +

^~^"^^

02 = FeQi,. (10)

The first reaction requires a knowledge of the

activity of iron in wiistite which can be obtained by

a Gibbs-Duhem integration of the oxygen activity.

Darken and Gurry, using Fe and O2 as components,

determined the activity of iron in wiistite from the

following relation:

Iogape = -l/2 r [(0/Fe-l)c?(logPoJ]

J 1.05

-l/2 1ogPo.+ l/2 logP;,,. (11)

They defined the activity of Fe at the Fe-FcjO

boundary to be unity, which is actually the standard

state for iron because of the low solubility of oxygen

in iron (2-3 ppm at 950 °C and 25 ppm at 1350 °C

[18]). f*02 is the partial pressure of oxygen corre- sponding to the two-phase Fe-Fe^O equilibrium.

^ The use of FeOy rather than Fe^'O is preferred since the O/Fe ratio is used in the Gibbs-Duhem integrations. 134
The second reaction described by eq (10) was

used in this study to determine the partial molal

quantities. In this case FeOi.os was selected as one

of the components and the formation of wiistite was

based on the solution of oxygen in the most oxygen

deficient wiistite phase, FeOi.os- It should be pointed

out that for a definite wiistite composition the sohd

solution under consideration is that of oxygen in the wiistite phase, not oxygen in iron. Thus, this selection of components appears to have a better

physical meaning than Fe and O2 except perhaps

for the use of magnetite and wiistite in describing

the binary solution. The composition of oxygendeficient wiistite does not vary appreciably with

temperature and can be considered constant over

the temperature range 1050 to 1200 °C [16]. The

standard state for oxygen was defined as an ideal gas at one atmosphere, and the activity of FeOi.os

at the oxygen-deficient composition was defined to be unity for aU temperatures. The standard free energy of formation and other thermodynamic

properties of the most oxygen-deficient wiistite (FeOi.05) were obtained from data in cell I [16]. The free energy of mixing of one mole of FeOi.05

with —— n^oles of O2 at any particular temperature is given by eq (12):

AF^u = RT\n apeoj.os + ^" ^^^^

From the knowledge of oxygen activity across the

variable composition of wiistite, the activity of

FeOi.05 can be determined by a Gibbs-Duhem inte- gration of the following equation:

lnapeo,05 = -r (y-l-05)c?(lnPo.)- (13)

Jl.05

Graphical integrations were performed on the iso- therms at 1057, 1075, 1100, 1139, 1150, and 1200 °C. The partial molal free energies, AT^peo, and

AFo, were calculated from the following relations:

^FeOi.o5 = -^FeOi.o5~-^FeOi.05""^^ ^^^601.05;

AFo = Fo-l/2F°, = ^ln Po.. (14)

The free energies for AFpeo, 05 are tabulated in table 2, and those for AFq can be calculated from

partial pressure of oxygen data contained in table 3. The toted free energy of formation of wiistite,

AF/(Fe0y)5 from the elements in their standard states of any composition y greater than 1.05 is given by

Table 2. Partial molal free energy of FeO, os in Wiistite

Composition 0/Fe 1.06 1.07 1.08 1.09 1.10 1.11 1.12 1.13 1.14 1.15

Isotherm Temp °C

Fed. 05

1057 6.6 21.2 41.8 64.2 91.5 136.5 190.0 227.0 242.0 285.0 1075 4.5 17.9 41.1 69.3 101.0 140.0 180.0 233.0 275.0 284.0 1100 Sample A, B, C 6.1 22.1 44.2 72.8 106.7 148.0 193.0 295.0 288.0 306.0 1139 7.0 21.5 37.7 62.5 103.2 156.4 217.0 273.0 326.0 344.0 1150 7.2 19.9 41.0 69.5 106.0 154.0 211.0 281.0 345.0 351.0 1200 8.2 25.0 51.9 85.0 124.3 172.5 225.0 278.0 333.0 368.0

'^FfM,M = fFeO, 05 - /^FeOl 05 = 2.303ft7" log Of^,

Table 3. Partial pressure of oxygen obtained from Po^ — OIFe isotherms

Composition 0/Fe 1.05 1.06 1.07 1.08 1.09 1.10 1.11 1.12 1.13 1.14 1.15

Isotherm Temp °C -log Po.

1057 13.93 13.58 13.26 12.98 12.78 12.58 12.32 12.01 11.87 11.81 11.79 1075 13.63 13.32 13.04 12.74 12.50 12.25 12.04 11.81 11.59 11.43 11.40 1100 Samples A, B, C 13.25 12.88 12.56 12.26 12.00 11.77 11.56 11.27 11.05 11.02 10.93 1100 Sample D 13.25 12.93 12.66 12.40 12.10 11.89 11.68 11.46 11.25 11.02 10.93 1139 12.73 12.32 12.00 11.80 11.62 11.29 10.98 10.73 10.47 10.28 10.23 1150 12.58 12.18 11.87 "11.61 " 11.35 " 10.10 10.88 10.64 " 10.39 " 10.18 10.12 1200 11.99 11.51 11.18 10.85 10.57 10.29 10.04 9.82 9.61 9.42 9.30

Average values between Samples A, D.

135
eq (15). AF/(peoi.o5) standard free energy of

formation,

^/(FeO;/)=^/(FeOl.05)+'^mixi (15)

of one mole of the most oxygen-deficient wiistite phase, FeOi.o5. It is known very accurately from

cell measurements of Fe-FcaO equilibrium de- scribed previously [16]. The values for free energy

of formation of FeOy are given in table 4. The partial molal enthalpies (A//o and A//peo, ,,.) were evaluated from a linear least-squares fit of the

log Oq and log aFeo, ,,5 data as a function of l/T. The

values of log P02' which were extracted from the

log Po2'0/Fe isotherms and used in all the thermodynamic calculations, are listed in table 3. The use

of a linear fit of log versus l/T to determine the

partial molal enthalpy of oxygen was justified on

the basis of the linear emf-temperature coefficients

obtained for different wiistite compositions. The

Table 4. Free ener

determination of the partial molal enthalpy of

FeOi.or, as a function of composition was also made

using a finear fit of log QFeo, 0.- versus 1/7. Although

there was no physical basis for a linear fit, the errors are small because the slopes of log a^eo versus \IT

plots are very small. The partial molal entropies were determined

from the temperature derivatives of the partial

molal free energies using a least-squares analysis

of the data. The partial molal enthalpies and entropies of oxygen at different compositions are given

in table 5 along with the least-square analysis of the

log^ P02 versus l/T data. A similar tabulation for

A//FeOi o5' '^FeOj ns' "FeOj 05 versus 1/r are found in table 6. The enthalpies of mixing were calculated from

equation (16) and are given in table 7:

A//mix= (y-1.05)A//o+A^Feo,.o5. (16)

offormation for FeOy

Composition 0/Fe 1.05 1.06 1.07 1.08 1.09 1.10 1.11 1.12 1.13 1.14 1.15

Isotherm Temp °C — (AF/(FeO,^) — AF/(FeO,.„5) (cal))

1057 44,447 44,865 45,275 45,675 46,070 46,455 46,835 47,195 47,565 47,930 48,325 1075 44,161 44.575 44,985 45,380 45,775 46,150 46,530 46.890 47,255 47,615 47,965 1100 Samples A, B, C 43,764 44,175 44,575 44,965 45,345 45,720 46,090 46,435 46,785 47,175 47,505 1139 43,144 43,550 43,940 44,325 44,710 45,070 45,428 45,785 46,125 46,465 46,795 1150 42,970 43,375 43,765 44,145 44,520 44.885 45,250 45,605 45,955 46,305 46,620 1200 42,175 42,470 42,955 43,325 43,685 44,035 44,380 44,715 45,045 45,370 45,680

1100 41.680 41,675 41,660 41,635 41,600 41,565 41,520 41,460 41,400 41,380 41,310 1200 40,165 40,065 40,145 40,115 40,075 40,030 39,980 39,925 39,865 39,800 39,720

Table 5. Partial molal enthalpy and entropy of oxygen as a function of wUstite composition

Composition

0/Fe

-log Po, = A + B X 10^

-(A//o)'' = -(AS„)'' =

-(so-|s°^)

A B cal call mole deg

1.06 -7.831 ±0.230 2.846 + 0.032 65.139+ 730 17.9

1.07 -8.371+ .491 2.879+ .068 65,862 + 1, .5.50 19.2

1.08 -8.875± .713 2.912 ± .091 66,197 + 2,080 20.3

1.09 -9.566± .914 2.975± .127 68,068 + 2.900 21.9

1.10 -10.512+ .573 3.071+ .080 70,2.53 ±1,830 24.0

1.11 - 10.924 ± .596 3.093+ .083 70,772 ±1,900 25.0

1.12 - 10.680 ± .722 3.024+ .100 69,200 ±2,290 24.4

1.13 -11.514+ .480 3.111+ .067 71, 180± 1,530 26.4

1.14 -12.844+ .583 3.275+ .081 74,927 ±1,850 29.4

1.15 -13. 690 ± .500 3.384+ .070 77,426 ±1,600 31.3

,,77 -2.303 X 1.987 „ ,,7; 2.303x 1.987 , "AZ/o^ Bx 10" and ASn = ^ A.

136
Table 6. Partial molal enthalpy and entropy of FeO,,o5 a function of composition

Composition

0/Fe log apuoi.os '4 + B 1 X 10^ \ LI C ^" FeOi.05 AC C '^^ FeOi.05

A B X lO-' cal caljmole deg

1 06 0.00397 ± 0.0019 - 0.409 ± 0.265 19. 0± 12.0 0.0182

1 07 .00594

±

.0043 - 0.363 ± 0.609 17. 0± 28.0 .0272

1 08 .01005

±

.0096 -0.463 + 1.33 21. 0± 61.0 .0460

1 09 .02080 + .0142 -1.36 ± 1.98 62.0+ 90 .0952

1 10 .03919 + .0115 -3.16 ± 1.61 144.0+ 73 .179

1.11 .05250

±

.00515 -4.01 ± 0.72 183.0 ± 32 .240

1 12 .06775

±

.0162 -5.00 ± 2.26 229.0 ± 103 .311 1.13 .0944 ± .0208 -7.56 ± 2.89 346.0+132 .432 1.14 .1505 ± .0403 -14.40 ± 5.0 660.0 ±228 .690

1.15 .1465 ± .0188 -13.35 ± 2.6 611.0+120 .670

<'A//reo,„j=-2.303x 1.987X 10<B and ASpeo 1.05= 2-303 x 1.987^.

Table 7. Enthalpy and entropy of mixing FeOj 0^ and oxygen and enthalpy and entropy of formation of FeOy

Composition

0/Fe 1 + y

A'S/lFeOj/)

caljmole caljmole caljg atom caljmole deg caljmole deg

1.05 0 65,609+ 50 32,005 0 15.13

1.06 636 66,245 ± 70 32,155 0.163 15.29

1.07 1,309 66,920+110 32,300 0.361 15.39

1.08 1,985 67,595 ±180 32,495 0.569 15.70

1.09 2,673 68,280 ±250 32,670 0.785 15.91

1.10 3,386 68,995 ±220 32,855 1.031 16.16

1.11 4,080 69,690 ±210 33,030 1.272 16.40

1.12 4,615 70,225 ±335 33,125 1.425 16.56

1.13 5,350 70 ,960 ±330 33,315 1.672 16.80

1.14 6,095 71,705 ±450 33,505 1.956 , 17.09

1.15 7.141 72.750 ±330 33,835 2.460

' 17.59

The total enthalpies of formation for any particular

composition of wiistite were calculated from equation (17) and are given in table 7:

^fif(FeOy)= ^H}^pgQ^^.^+ ^.HjrMx- (17)

The entropies of mixing and formation as a function of composition were calculated in a similar manner and are tabulated in table 7. Additional thermodynamic data were obtained by

performing emf-temperature profiles for successively titrated compositions. Using cell IV, a number

of oxygen-pressure versus temperature profiles were obtained between 700 and 1150 °C at a constant heating or cooling rate which was varied

between 0.30 and 1.25 °C/min. Within the wiistite single-phase region it was difficult to retrace the

emf profiles upon heating. This problem was probably due to variations in gas atmosphere as a result of outgassing effects upon heating. Thus, for most

of the profiles taken, only the cooling curves were

used. A typical profile is shown in figure 7. Note

that upon heating the increase in voltage indicates

that the sample is being oxidized. It was also observed that breaks in the emf-temperature curves occurred between 800 and 1000 °C and did not appear to be a function of the heating or cooling

rates. However, there was considerable hysteresis

in the temperature break and the slopes in the lower

temperature regions. The slopes for the high temperature regions were reproducible. A linear least- square fit of log f*02 versus l/T was made on each

profile in selected temperature regions where a break in the curve occurred. The least-square

analysis (see table 8) confirmed the existence of more than one slope.

Because of the problem of the oxidizing-reducing

cell atmosphere it was difficult to establish the

actual composition. An alternate method of deter- mining the composition for a particular emf-temperature profile was to match the partial pressure

of oxygen with each of the oxygen pressure-compofeition isotherms (1057 to 1200 °C). For example,

the logarithm of the partial pressure for composition

"C" at 1057 °C determined from -log Po2 = - 7.206

+ 2.777 (l/r)XlO^ (table 8) was -13.674. The

composition (0/Fe) corresponding to this pressure on the 1057 °C isotherm is Fed.0.57.5. This procedure

290-621 0-68— 10 137
Figure 7. Temperature profile of cell V for composition "E".

was repeated for each isotherm and the resuhs are recorded in table 9. The high-temperature isotherms were selected in order to eliminate any ambiguity

due to changes in slopes or possible nonequilibrium conditions that occurred below 1000 °C. It was

noted that the compositions determined from lowtemperature isotherms (800 and 894 °C) did not agree with the determined average composition.

The partial molal enthalpies (Hq— 1I2HqJ were

determined from least-squares analysis of log Pq^ versus IIT and are presented in table 8. The partial

molal entropies were determined from the inter- cept of the equation for log Pq^, which agreed well

with the values obtained from the emf-temperature

coefficients.

The partial molal enthalpies and entropies,

A//o and ASq, for these compositions are plotted

with other data obtained from the isothermal ti- tration experiments in figure 8. Although the error

limits plotted with the data from the titration experiments are rather large compared with these data,

the two sets of data appear to follow the same trend. That is, the partial molal enthalpy for oxygen levels

off at a composition between FeOi.io to FeOi.13

before again becoming more negative. This change

in slope of A/Zq could be indicative of a change in the defect structure of wiistite. This change in slope occurs in the composition range in which a positive

deviation from Raoult's Law was noted. The in- flection point of P02 versus 0/Fe also corresponds to

Table 8. Linear least-squares fit of log P02 in equilibrium with FeOy compositions determined from cell V

Composition ^ Temp range. -logPo.=^ + B(^)xlO^

So-l/2S; A B

°C caljmole caljmole deg "B" 850-1040 - 7.0725 ± 0.0083 2.7763 ± 0.0010 63,522+ 25 -16.18

"C" 850-1125 - 7.2061 ± .0720 2.7774± .0094 63,547 + 215 - 16.49 "D" 825-1150 -7.845 ± .020 2.8224

±

.0025 64,577 ± 58 - 17.94 "M" 950-1000 -9.187 ± .10 2.92 ± .02 66,810 + 460 -21.00

800-900 -10.966 ± .105 3. 1294 ± .0118 71,601 ±270

800-1000 -10.019 ± .180 3.0427 + .02106 69,617 + 480 "E" 875-1100 - 8.7394 ± .0110 2.8555 ± .0014 65,334+ 30 -19.97

750-875 - 9.6092

±

.0360 2.9559 ± .0039 67,631+890 "K" 750-1100 -9.054 ± .060 2.8956 ± .0071 66,251 + 160 "F" 925-1100 - 9.8815 ± .0346 2.9623 ± .0044 67,777 + 100 -22.61

800-900 - 9.7042 ± .0630 2.9403 ± .0071 67,274 ±160

800-1100 - 9.7704

±

.0238 2.9479 ± .0029 67,450 ± 65 "G" 750-1100 -9.840 ± .013 2.9188 + .0017 66,782 ± 40 -22.45 "H" 1000-1200 -8.997 ± .071 2.7989

±

.0098 64,039 ± 220 -20.40

"1" 950-1050 - 12.46 ± .1 3.23 ± .01 73,904 + 230 -28.51

800-900 -11.992 ± .059 3.1792 ± .0066 72,740 ±150

800-1050 -12.440 ± .064 3.2294 ± .0076 73,889 ±175

"J" 1000-1150 -12.791 ± .106 3.2619± .0143 74,632 ± 325 -29.2 "K" 950-1100 -12.826 ± .021 3.2653

±

.0027 74,710 ± 60 -29.3

775-975 -11.847 ± .047 3.145 ± .0054 71,958 ±125

775-1100 -12.222 ± .079 3.187 ± .0095 72,919 ±215

^ See Table 9 for actual compositions. 138
Table 9. Composition of w'ustite derived from comparison of P02 ivilh various Po^-FeOy isotherms

Temp. °C of Po.-0/Fe Isotherm

Composition Designation

B C D E F G H I J K L

1057

1075

11 no

1139 1150

1200

0/Fe Ratio'

1.054

1.054

1.0528

1.0535 1.0535

1.0575

1.058

1.058 1.056

1.058

1.057

1.067 1.069

1.067

1.066

1.066

1.067

1.091

1.091 1.089

1.094 1.093

1.091

1.108 1.108

'^1.106

1.118

1.121 n.U9

1.120 1.124

« 1.120

1.136 1.136

" 1.134

1.163 1.146 "1.138

1.138

" 1.152

1.145

1.163

1.143

« 1.138

Avg 0/Fe Ratio 1.0536 1.058 1.067 1.090 1.093 1.107 1.119 1.121 1.135 1.148 1.148

'0/Fe ratio determined from matching Pq, of each composition (C. D. etc.) to that of Po,OIFe isotherms.

" Average values based on more than one sample.

the composition where wiistite changes from a

p- to n-type semiconductor [19-21]

.

The resuhs o.f Darken and Gurry [9] at 1250 °C

show that A//o varies from —61.72 kcal/mole

(FeOi.055) to —63.78 kcal/mole for FeOi.n compared

62

AHQ-Determined from Single Wustite Compositions

AHQ"Defermined from Titration Data

1

64 \

Q

1 N

66 1

1

68 1

\ \

\

70 \

X

72 0

\

\ 0

\ \

74 •

\

\

\

N

V\ 76 0 • ASq "Determined from Titration Dota 0 ^ " -Q 0 ASq -Determined from Single Wustite Compositions •

7R 1

1

1 1

12

<] 20 '

24

28

I I I I I 1-05 107 109 ill 113 1-15 0/Fe

Figure 8. Partial molal quantities AH,„ AS,, /or wiistite.

to —64.5 and —74.5 kcal/mole in this study. This

comparison may not be valid since the data reported

here extend from 1057 to 1200 °C compared with

the 1100 to 1400 °C range studied by Darken and

Gurry. The thermodynamic properties of wiistite determined by Kleman, Raccah, and VaUet [12, 22] were based on the presence of three wiistite phases,

and are described by equation (18):

log Po, = {a/r+ 6) (0/Fe) + c/r+c^. (18)

The coefficients (a, b, c, d) were defined for each of the three wiistite phases proposed. Kleman

shewed how the results of Darken and Gurry for A//o could be duplicated by the average partial

molal enthalpy for each of the three phases. The

proposed phase diagram and plot of partial molal

enthalpies of oxygen for the three phases are shown

in figure 9. Note that wiistite III is stable Jjelow

1010 °C. The partial molal enthalpies ^Ho for

wiistite II ( — 65 kcal/mole Fed.05, ~67 kcal/mole,

FeOi.14) are similar in magnitude to those reported

in this study. However, for wiistite I, which should

correspond to data of Darken and Gurry (1100 to 1400 °C), A^o varies from -68_kcal (FeOi.0.5) to — 60 kcal/mole (Fed. 14)- Other A/Zq values reported

vary between -66.5 to -69.6 kcal [20, 33, 34]. The

partial molal enthalpies of oxygen determined by

two different methods in this study were consistent

in calculating higher values than those reported by

Darken and Gurry.

A comparison of the enthalpy of formation as a function of composition with Darken and Gurry's

results at 1250 °C indicated agreement within 500

cal. It has been proposed by Ariya and Morozova

[33] that nonstoichiometric compounds can be con- sidered to be solid solutions of one stoichiometric

139
compound in a second stoichiometric compound.

Thus, considering FeOi.os and Fe304 as the two

components, the enthalpy of formation per gram

atom (see table 7) as a function of composition does

indeed He very close to the extrapolated line drawn

between Fed.05 and Fe304.

Discussion

The question of whether diffusion is fast enough

in the oxygen-rich compositions to eliminate polarization of the cell causing non-equilibrium conditions

must be considered. It has been proposed by

Raleigh [24] that diffusivities lower than 10~^

cm-/sec can lead to excessive times for the elimination of concentration gradients. Diffusion constants between 2 X 10"'' and 1 X 10"'^ cm-/sec have been

reported for iron in wiistite at 1000 °C. [25-28].

Most of the reported data indicates that the diffusion constant for iron in wiistite decreases with increasing oxygen content and, that in magnetite, it is two

orders of magnitude lower than in FeOi.os- Thus,

assuming that diffusion of iron was rate controlling,

then during the titration of oxygen into FeOy,

polarization should occur with the following relation

holding:

40=[^4Fe+^

dt >dt

(19)

With the rate of arrival of ions at the ZrO^-FeOy

interface greater than the transport of Fe-+ to this interface, the oxygen potential would change rapidly

and, indeed, a blocking layer of Fe304 would result eventually and be recorded by a marked change in

cell voltage. It is noteworthy that the presence of

any Fe304 at the zirconia interface would result in 400 to 500 mv change in cell voltage. Using the same

logic an attempt was made to oversaturate the inter- face by titration of oxygen for the purpose of measuring the mobility of holes or the predominant

conducting species in wiistite. However, titration of 11.8 ma into FeOy for 75 min. at 1084 °C at a current density of 12.5 ma/cm- did not result in any polarization. Assuming a mobility of 0.11 cm''^/ volt sec for holes and with the number of vacancies

equal to number of holes, the current calculated

with the known dimensions and emf across the

FeOy electrode, required to achieve this mobility, was approximately three orders of magnitude

greater than 11.8 ma. An experiment which also provides information on the diffusion of iron in the FeO system was

performed by Barbi [29]. The purpose of his experiment was to determine the free energy of formation

of AI2O3 from the following cell:

1400

1000

600

I 04

Figure 9a. Fe-0 phase diagram after Kleman [12].

9b. AH„ as a function of different wiistite, after Kleman [12}.

Current was passed through the anode resulting in reduction of the AI2O3 film on the Al metal, and

oxygen entering the cathode oxidized the Fe-Fe^O

to Fe203. Measurement of cell emf versus time

after titration resulted in two regions of constant

cell voltage. The constant emf indicates that a two-phase region is being reduced. The difference

in cell voltage between these two constant-emf

regions corresponded to the change in oxygen

potential between the two-phase equilibria Fe203-

Fe304 and Fe-FcxO. A typical time required for reduction from Fe203 back to the Fe-Fe^O mixture was 10 minutes for a titration current of 3 ma for 25 minutes at 639 °C. This rapid reduction cer- tainly indicates that equihbrium is achieved rapidly

in the Fe-0 system.

Another question to consider is the effect of impurities on defect phases noted. The purity of

iron was examined by starting with a sample of

pure zone-refined iron (< 25 ppm impurities) and

Pt, Graphite, A1+ AI2O3 Th02

15 mole % LaO,.5

Fe + Fcj-O, graphite, Pt.

140
titrating oxygen into the iron until the Fe-Fej-0

boundary composition was exceeded (Sample

C — cell III). Titration was performed at temperatures between 900 and 1000 °C in order to reduce

any effect due to reaction between platinum and

the iron sample. Assuming a composition of FeOi.05

as the iron-rich boundary, the error in the number

of coulombs titrated to this boundary was 0.9 per- cent, which is well within weighing errors for the

92.2 mg sample of iron. The results of the 1100 °C

isotherm are shown in figure 4. The only discernible

difference was a 3 mv higher cell potential at the

high oxygen compositions compared to other experiments at the same temperature on different samples.

The results of this study support and perhaps can be used to explain various observed phenomena

in the iron-oxygen system. Consider the transition

in wiistite from a p- to n-type semiconductor. The

arrows in figures 3-6 indicate the pressures [20] where the p-n transition occurs. It has been reported

that the conductivity becomes insensitive. to oxygen

pressure beyond the transition point and the

composition of this transition decreases with in- creasing temperature [20]. The transition appears

to correspond to the inflection point of the curves which indicates a change in the activity-composition

coefficient. The change of the coefficient is cer- tainly suggestive of a change in defect structure.

It is noteworthy that Kleman [12] apparently

interpreted this effect as a phase change within

the wiistite region. The results of this study between

950 and 1200 °C do not substantiate his proposed

diagram (fig. 9). The presence of any two-phase

field within his diagram would certainly have been

observed with the sensitive coulometric titration technique employed in this study. It is easy to understand how such an interpretation could be

made, based on plots of log P02 versus 0/Fe ratio which do show changes in slope.

A plot of log P02 versus 1/7", corresponding to the

oxygen pressure at which p-n transition occurred,

showed two distinct slopes. The partial molal

enthalpy of oxygen was 67.0 kcal/mole between

950 and 1100 °C and 45.8 kcal/mole between 1100

and 1300 °C. It is noteworthy that at lower temperatures the wiistite compositions where the transition occurs are FeOi.io to FeOi.n, and that the 67

kcal/mole corresponds to the partial molal enthalpy

of oxygen determined for this composition range. Libowitz [30] has recently derived a model

expression relating the activity of oxygen (P^-)

to the deviation in stoichiometry 8 in FeOi + g. The

model is based on the assumption that the principal

defect is a complex composed of two iron vacancies

associated with an Fe+'' interstitial as proposed

by Roth [31]. The model equation was applied to

the isotherms in this study. It was found that the ex- pression predicted the Pq^ versus 0/Fe dependence

fairly well between Fed.05 and FeOi.n and at temperatures between 1057 and 1200 °C. The correlation

between the model and the titration data provided

additional support for the selection of the (1057 to 1200 °C) temperature range in which to determine

the thermodynamic functions of wiistite. Below

this range the correlation was unsatisfactory. It

is interesting to note at this point that Koch [32] has recently observed by high-temperature x-ray

diffraction that vacancies form clusters around iron

ions, presumably Fe+^ ions, and that reflections

exhibit characteristics of superlattice lines indicat- ing ordering.

Conclusions

The usefulness of a solid electrolyte galvanic

cell in determining phase boundaries and thermodynamic functions in the iron-oxygen system has

been demonstrated. Coulometric titration of oxygen

through calcia-stabilized zirconia makes it possible

to explore the entire iron-oxygen system with one

experiment.

The activity of oxygen as a function of wiistite composition was determined at temperatures be- tween 1057 and 1200 °C. The results led to the de- termination of partial molal quantities for the

solution of oxygen in FeOi.05 for all wiistite compositions FeOy. It is noteworthy that the partial

molal enthalpy of oxygen appeared to remain constant between FeOi.io and Fed. 13 before again

becoming more negative.

A statistical mechanical model, developed by

Libowitz, based on Fe+' interstitials associated

with two Fe vacancies, predicted correctly the

P02 versus 0/Fe dependence at temperatures be- tween 1000 and 1200 °C.

The writers are indebted to the Research Committee of the University of Utah for providing funds

for the research. The principal author is grateful

to Brig. General Raymond A. Gilbert and the Air

Force for affording the opportunity to continue

studies in the Ph.D. program.

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[14] G. L. Humphrey, E. G. King, and K. K. Kelley, U.S. Bur.

of Mines Rep. Invest., No. 4870 (1952).

[15] C. E. Wicks and F. E. Block, Thermodynamic Properties

of 65 Elements — Their Oxides, Halides, Carbides, and

Nitrides. BuU. 65, U.S. Bur. Mines, pp. 57-59 (1963).

[16] H. F. Rizzo, R. S. Gordon, I. B. Cutler, The Determination

of Phase Boundaries and Related Thermodynamic Properties in the Fe-0 System by Coulometric Titration, to be

published.

[17] H. F. Rizzo, R. S. Gordon, and I. B. Cutler, Evidence for the Existence of More Than One Defect Phase Between

Wiistite and Magnetite, unpublished work to be pub- hshed.

[18] J. H. Swisher, and E. T. Turkdogen, Solubility Permeability, and Diffusivity of Oxygen in Solid Iron, Trans. AIME, 2 2 9 , 426-31 (1967).

[19] I. Bransky and D. W. Tannhauser, High Temperature Defect

Structure of Ferrous Oxide, Trans, AIME, 2 39 , 75-80

(1967).

[20] G. H. Geiger, R. L. Levin, and J. B. Wagner, Jr., Studies on the Defect Structure of Wiistite Using Electrical Conductivity and Thermoelectric Measurements, J. Phys.

Chem. Solids, 27, 947-956 (1966).

[21] D. S. Tannhauser, Conductivity in Iron Oxides, J. Phys.

Chem. Solids, 23, 25-34 (1962).

[22] Pierrie VaUet and Paul Raccah, Contribution of the Study

of the Thermodynamic Properties of Ferrous Oxide, Mem.

Scien. Rev. Metall., 62 (1), 1-29 (1965).

[23] S. M. Ariya and M. P. Morozova, E. Vol'f, The Chemistry

of Compounds of a Variable Composition (English Translation), Zh. Neorgan. Khim U.S.S.R. 11, 13-22 (1957).

[24] D. O. Raleigh, Solid State Electrochemistry, Progress in Solid State Chemistry, 3, 83-134 (1966).

[25] P. F. J. Landler and K. L. Komarek, Reduction of Wiistite Within the Wiistite Phase in H2-H2O Mixtures, Trans. AIME, 2 36, 138-149 (1966).

[26] L. Himmel, R. F. Mehl and C. E. Birchenall, Self-Diffusion of Iron in Iron Oxides and the Wagner Theory of Oxidation. Trans. AIME, 197, 827-843 (1953).

[27] R. L. Levin, J. B. Wagner, Jr., Reduction of Undoped and Chromium-Doped Wiistite in Carbon Monoxide-Carbon

Dioxide Mixtures, Trans. AIME, 233, 159-167 (1965).

[28] C. T. Fujii and R. A. Meussner, Determination of Chemical

Diffusion Coefficient of Iron in Wiistite at 1000 °C, Report of NRL Progress, PB 174294, Naval Research Laboratory, Washington, D.C., pp. 27-28, March 1967.

[29] Giovanni B. Barbi, Thermodynamic Functions by E. M. F. Measurements on Solid Galvanic Cells in Non-Stationary

Conditions, Systems Al-fAl203, Trans. Faraday Soc,

62, 1989-1595 (1966).

[30] G. G. Libowitz, Defect Complexes and Microdomains in Oxides, paper presented at this conference.

[31] W. L. Roth, Defects in the Crystal and Magnetic Structures

of Ferrous Oxide, Acta. Cryst., 13, 140-149 (1960).

[32] F. Koch, Private Communication.

[33] S. J. Ariya, M. P. Morozova, and A. L. Shneider, I. The

Thermodynamics of Ferrous Oxide, J. General Chem.

U.S.S.R., 24, 37-42 (1954).

[34] L. Himmel, Self Diffusion of Iron in Iron Oxides and the Wagner Theory of Oxidation, unpublished Ph. D. Thesis,

Carnegie Institute of Technology, Pittsburgh, Pa. (1953).

142
The Mechanism of Oxygen Self-Diffusion in

Nickel and Cobalt Oxides

M. Hoch and R. Szwarc

Materials Science Program

University of Cincinnati

Cincinnati, Ohio 45221

The cation diffusion coefficients and hole conductivity data in FeO, CoO, and NiO were correlated

assuming a vacancy mechanism for cation diffusion

and a hopping mechanism for hole diffusion. The

cation diffusion coefficients in Fci-gO, Coi-sO,

and Nii-gO were shown to be proportional to the

concentration of vacancies 8 in the cation lattice with identical pre-exponential and exponential

terms. The hole diffusion coefficients have the same pre-exponential term as that for cation

diffusion, but the activation energy depends on

the bonding of the hole to the respective ion. Only

one of the two holes created by the introduction of a cation vacancy takes place in the conduction

process. The sintering kinetics of pure and doped nickel

and cobalt oxides were studied in an effort to elucidate the mechanism of oxygen self-diffusion

in these oxides. Both initial and intermediate

sintering models were used to study the oxygen

diffusion coefficients. Doping with monovalent

lithium ions increased the oxygen diffusion coeffi- cient. Addition of trivalent chromium ions de- creased the oxygen diffusion coefficient, although

the effect was not as great as in the case of lithium.

The defect responsible for oxygen self-diffusion in CoO and NiO is thus a complex, made up of an oxygen vacancy and two electron holes (Fo0©).

This model was used to calculate the Schottky

product in NiO and CoO; they were found to be

equal.

This research has been supported in part by the

Air Force Materials Laboratory, Research and

Technology Division, Air Force Systems Command,

United States Air Force, under Contract No.

AF 33(615)-2746.

143

Experimental Evidence for Highly Mobile Electrons in MnO and

NiO at High Temperature

D. S. Tannhauser,^ N. M. Tallan - and M. Gvishi

Department of Physics

Technion — Israel Institute of Technology, Haifa, Israel.

The measured Hall mobility in MnO and NiO

near 1000 °C, is n-type in the range where the

conductivity is p-type. For MnO, where the measure- ments extend into the n-type conductivity range,

this gives directly /i^'=10 cm-/F-sec. For NiO

the analysis is indirect, it rehes on a model with

neghgible HaU mobihty of the holes, and on the

experimental value of the optical band gap. The

A^lue calculated is [jl^ ^ 30 cm^/V sec. The drift mobiUty of holes in the compounds

MnO, FeO, CoO and NiO is about 0.1-1 cm^/F • sec

at r ^ 1000 °C. This low mobihty is connected with

the fact that the holes move in the c?-levels of the

cations, however the question whether the motion

is by band conduction or by small polaron hopping

is not settled. We would hke to present experimental

evidence that in MnO and in NiO the mobility of

the electrons is much larger than that of the holes. The evidence for MnO is much more direct than that

for NiO.

Hed and Tannhauser [1] ^ combined conductivity with gravimetric measurements on MnO

at r^llOO °C. The partial pressure of oxygen

in the atmosphere surrounding the sample (estabhshed by mixing CO2 and CO in measured ratios 9 = P(C02)/P(CO)) controlled the deviation from

stoichiometry and therefore also the concentration of holes and electrons. It was found that the

conductivity goes through a minimum at q—l,

holes dominate the conductivity at g > 1 and elec- trons at g < 1. The detailed analysis of the conductivity and gravimetric data gave )Mp = 0.2 crn^/V • sec

at 1200 "C, and jJinll^p ^ 10 ^t the same temperature;

the lower hmit reflecting the fact that the weight

change in the n-type region was immeasurably small.

ix-n and /i,p are the drift mobihties of electrons and

holes. Thermoelectric measurements in ref [1] and similar data of McKinzie [2] confirmed that

f^nll^p > 1? but the numerical value of this ratio

' At present at the Division of Engineering, Brown University, Providence, R.I. 02912,

- On leave from Aerospace Research Laboratories, Wright-Patterson Air Force

Base, Ohio 45433.

'^Figures in brackets indicate the literature references at the end of this paper.

could not be accurately determined because of the uncertain values of the heat of transport.

Recent Hall effect measurements gave a ratio of

l^nlfJ^p ^ 50 at 943 °C, where fx" are the Hall mobilities. An a-c Hall effect apparatus with a 510 cps

electric current and a 2 cps magnetic field was

used [3]. The data (see fig. 1) were taken on a single

crystal of MnO in which the charge carrier concentrations were varied by a factor of 30 through control

of PiO-z), spanning the range from electron controlled to hole controlled conductivity. The Hall

effect stayed n-type in all the range where we could measure. Figure 1 shows that the dependence of the measured ^tHaii on PiOz) agrees with the theory

[MHaii = )^n(l-a/6)/(l + «) where a = p/x'^lnfji^ and

b = iJ^nlfJ''"] for a two carrier semiconductor with

IJb^, = 10 cm'IV sec and fM^I/Ji^ = 50. For the highly

mobile electrons it is fair to assume that fi" =^t^

and furthermore the Hall mobility of the holes,

which is 0.2 cm'/F-sec, agrees fairly well with

their drift mobihty extrapolated from the data of ref [1]. It is therefore likely that the value of fJi^lfip at 943 °C is also 50, in agreement with the estimate

of ref. [1]. ix^ could not be measured directly be- cause of electrical noise generated by the contacts

at high deviations from stoichiometry.

In NiO similar measurements were made at 1007 °C. In this compound, similarly to MnO, the

Hall effect was n-type, while the conductivity was

clearly p-type both by its dependence on PiO-z) and

by thermoelectric data over the whole range of

measurement.

The relation p-Haii = Ccr~'~ was found to fit the

data over the range of measurement (see fig. 2). An explanation based on the two carrier model

can be given, however, here the relation /i,{j = />t^ does not hold. It is known from conductivity vs. P(02) experiments that a> I and therefore a = K(T'^ so that the above equation gives

/^Haii = ^n(l/« - 1/6) = Kix^a-'- \x»

and this fits the experimental curve provided

145
JO'

10"

q = PCCOgVPCCO)

Figure 1. The Hall mobility /XHaii of MnO at 943 °C plotted against q =P(CO,)/P(CO).

® are experimental points, the lines are theoretical curves drawn with the assumption /i."= 10 cm-/r sec and various values of b = fj."lfi". Values of a = q'^'^ are also given on

the abscissa.

"

CP 10

10"

T—I II I

I \

\

\

\

\ \

\

\

J I

I I I I I J I I I I I 10" 2 5 CT (ir'-cm"')

Figure 2. The Hall mobility /XHaii of NiO at 1007 °C plotted

against the conductivity u.

The dashed line represents the relation y.^^^y = C<j~~. The point at o" = 3xl0~'.

which was taken in air, represents an upper limit only.

eV) gives /Lt„ > 30 cm^jV • sec, depending on whether

the effective density of states for electrons is taken

as A^e= 1 per ion-pair or A^e < 1- It is a priori reasonable that the mobility of

electrons in transition metal oxides is much larger

that that of the holes, since electrons are pre- sumably located in the 45 states of the cations which

have considerably larger overlap than the 3d states. The high values of /jl,, in MnO and NiO support

this picture and it is therefore reasonable to assume

that electrons move by band conduction in these

oxides.

There is also experimental evidence from thermoelectric data that electrons play a role in CoO,

however no measurements of the electron mobility

are available.

fjL^ < /AHaii at the low (T end of the measurement

range. Therefore yu,^ <^ 2.10"^ cm'-/F • sec while

(JL^ = 0.5 cm'-IV • sec. The relation ijl^, < fJifj might be

explained by a recent theory by Maranzana [4] which has already been used by van Daal and Bos- man [5] to explain their results for Li-doped NiO;

alternatively it might point to hole conduction by

small polarons where the ratio fiplfJ.^ depends very

sensitively on the details of the model chosen.

The electron mobility cannot be directly cal- culated in the case of NiO, the experimental data

giving only the value of the product nfji,i. Estimating

the product n p from the optical band gap (A£' = 4

This work was partiaUy sponsored by the Aero- space Research Laboratories under Contract No.

AF 61(052 )-825 with the European Office of Aero- space Research, U.S. Air Force.

References

[1] A. Z. Hed and D. S. Tannhauser, J. Chem. Phys. 47, 2090

(1967).

[2] H. McKinzie, Ph. D. Thesis, Arizona State University (1967).

[3] N. Z. Lupu, N. M. TaUan and D. S. Tannhauser, Rev. Sc.

Instr. 38, 1658 (1957).

[4] F. Maranzana, Phys. Rev. 160,421 (1967).

[5] H. J. van Daal and A. J. Bosman, Phys. Rev. 158, 736 (1967).

146
High Temperature Defect Structure and Electrical Properties of NiO

I. Bransky

Israel Institute of Technology

and

N. M. Tallan

Aerospace Research Laboratories

Wright-Patterson Air Force Base, Ohio 45433

Although there is general agreement that pure

NiO is an oxygen excess semiconductor ^roughout

its stabiUty range and that the predominant defects

are metal vacancies and holes, both singly ionized

[1, 2]' and fully ionized [3] vacancies have been

proposed. There has been a similar lack of agree- ment concerning the nature of the charge carrier

transport mechanism. Although most of the early

investigations on Li-doped NiO concluded that the

charge carrier transport in NiO is thermally acti- vated, as expected for small polarons in the hopping

regime, several recent investigations have presented

evidence for a narrow band conduction mechanism

at temperatures below about 1000 °C [4, 5]. The objective of this study was the determination of the temperature and oxygen partial pressure

dependence of the electrical conductivity and

Seebeck coefficient of NiO above 1000 °C. Both

single crystal and polycrystaUine specimens were

studied. The single crystal specimens were epitaxially grown on MgO crystal substrates by the

technique of vapor decomposition of NiBr2. The

polycrystaUine samples were prepared by oxidation of spectroscopically pure nickel foil. The various oxygen partial pressures required were estabhshed using pure oxygen, oxygen-argon

mixtures, or CO-CO2 mixtures. Four terminal

conductivity measurements were made to eliminate contact resistances. An a-c bridge operating at 80 cps was used after prehminary measurements

showed no frequency dependence of the conductivity up to 20 Kcps. The thermal emf measure- ments were made by heating one end of the

specimen to estabUsh a steady state temperature

gradient in it, and then measuring the voltage

developed between the pure platinum legs of two

thermocouples, one at each end of the specimen.

' Figures in brackets indicate the literature references at the end of this paper.

Although in all cases the voltage developed was

significantly larger than the residual voltage at zero temperature gradient, the thermal emf values

reported were obtained by least-square analysis

of many determinations at different temperature

gradients.

The time required for a 0.33 mm thick single

crystal specimen to attain steady-state conduction

values, after isothermal changes in the ambient

oxygen partial pressure, was determined as a pre- liminary step in the conductivity studies. Above

1200 °C the apparent equilibration was accomplished within ten minutes; below 1200 °C, considerably longer times had to be allowed to obtain

reproducible, meaningful results. The oxygen

partial pressure dependences of the conductivity

of single crystal and polycrystaUine specimens at constant temperature were of the form cr °c /'(Oa)"'^, where x varied from 4.1 ±0.1 at 1000 °C to 4.3 ±0.1 at 1600 °C. This pressure dependence is consistent with the defect formation reaction

¥2 O2 (gas) ^ Oo-\-Vm +®^ describing the formation of single ionized metal vacancies and holes. The temperature dependence of the single crystal

and polycrystaUine specimen conductivities were

identical between 800 and 1500 °C. Assuming a pre-exponential dependence in the conductivity,

and therefore plotting log aT versus T~^, an

activation energy for conduction of 1.03 ±0.02 eV

was found. Since the observed pressure dependence

of the conductivity suggests the formation of singly

ionized metal vacancies and holes, this activation

energy for conduction includes the energy of formation of the defects, A////2, and any motional energy

of the hole. The Seebeck coefficient, a = emf/Ar, was measured against platinum on a single crystal specimen.

AU Seebeck coefficient determinations made were

p-type, indicating a predominance of positive

147
charge carriers. A value for a. was obtained at each

temperature studied by a least squares analysis

of at least twenty measurements of the voltage

developed across the specimen as a function of

the temperature gradient imposed. The Seebeck

coefficient was found to be exponentially temperature dependent with an activation energy of

0.66 ± 0.03 eV. In an extrinsic p-type semiconductor, a = A/e (In Ni^jp + A), where A^^ is the density of

states and A is the kinetic energy transported by

the charge carrier. The energy of formation of the

defects, A////2, is therefore 0.66 eV if A^;, is temperature independent, as generally assumed for a

localized carrier model, and 0.75 eV if A^i varies as

T^l-, as in the case of a broad-band semiconductor.

These values for A///72 are in reasonable agreement

with values of 1.36 ±0.04 eV [2] and 1.5 + 0.2 eV

[3] for AHf obtained from conductivity measure- ments alone. The activation energy for conduction 1.03 eV,

containing both A///72 and any energy required for the motion of the carriers, is therefore significantly

larger than the value of A////2 found from Seebeck

coefficient measurements at high temperatures.

This is in direct support of the appHcability of a localized model for the charge carrier and a hopping

model for its transport mechanism at high temperatures. Subtracting the value 0.66 eV for A////2 from

the energy for conduction, a value of 0.37 eV for the motional energy of the hole is obtained.

References

[1] H. H. V. Baumbach and C. Wagner, Z. Physik. Chem. B24,

59 (1934).

[2] N. G. Eror, Thesis, Northwestern Univ., 1965.

[3] S. P. Mitoff, J. Phys. Chem. 35, 882 (1961).

[4] A. J. Bosman and C. Crevecoeur, Phys. Rev. 144, 763 (1966).

[5] I. G. Austin, A. J. Springthorpe, B. A. Smith and C. E. Turner,

Proc. Phys. Soc. (London) 90, 157 (1967).

148
Ionic Conductivity in Oxides:

Experimental Problems

Survey of Existing Data

L. Heyne

Philips' Research Laboratories

Eindhoven, Holland

Electrically insulating oxides may show electronic as well as ionic conduction when heated to a high temperature. To characterize this high temperature conductivity, it is necessary to separate both contributions and to study their dependence on oxygen pressure and doping. The theory is dis- cussed on which methods for determination of the conductivity components are based. Starting from

general charge transport equations, it is possible to express electronic and ionic currents in terms of two gradients; that of the electrochemical potential of electrons, and that of the chemical potential of oxygen. The first is measured with metal contacts, the second determined by the local oxygen partial pressure. The basic assumption is made, that local thermodynamic equilibrium between all charge

carriers and the lattice is undisturbed by current flow. In practice, different boundary conditions can

be obtained by a choice of electrodes with special properties with respect to reversibility or blocking

action. This opens the possibility of the determination of transport numbers. Four different situations

will be considered: both contacts reversible and in the same atmosphere: both contacts reversible but at different oxygen pressures; blocking of the ion current; blocking of the electron current. Special

attention is paid to the question of how close the situation on which the quantitative evaluation of the experiment is based is approximated by the practical arrangement. Finally, a short review is given of the most important properties of those oxides, about which dependable quantitative information is available in literature.

Key Words: Blocking electrodes, electronic conductivity, ionic conductivity, reversible electrodes, thermodynamic equilibrium, transport numbers.

1. Introduction

The electrical conductivity of solids can be a consequence of the presence of mobile electronic

charge carriers as well as of ionic charge carriers. In the general case both types of conduction are present simultaneously. The total electrical conductivity must then be characterized by the partial

conductivities of electrons and holes and of the

different ionic species that can be present. In an

important respect mixed ionic and electronic semiconductors differ essentially from purely electronic

semiconductors. This is a consequence of the fact

that, since ionic conduction forms an integral part

of the total conductivity, the substance can come

easily into equilibrium with the surrounding at- mosphere. This means that significant values of

conduction parameters can only be determined if due care is taken to control the composition of

the sample. This can be done by controlling the

surrounding gas atmosphere, or by equilibrating

the sample with a solid in which the thermodynamic

potential of one of the components of the sample

is fixed (for a binary substance). This solid can

often be chosen as one of the electrical contacts. In this survey attention will be concentrated

only on those aspects of mixed conduction that have

a direct bearing on the methods of determination

of the various components of the conductivity.

The description of ionic conductivity in terms of

fundamental lattice properties and lattice defects

will not be undertaken here [1].

' A complete characterization of the electrical conductivity of an oxide consists of the determination of the electronic and ionic partial conductivities

as a function of the oxygen partial pressure (or the

deviation from stoichiometry) and of the concentration of added impurities. A knowledge of these

dependences can lead to a quantitative description

of the solid under consideration in terms of the

thermodynamics of lattice defects. Here, only the ionic conductivity caused by ionic

species of the compound itself will be considered.

This means that it will be assumed that the mobility

of foreign ions is essentially zero and, therefore,

that the only influence of their possible presence

' Figures in brackets indicate the literature references at the end of this paper.

149
is on the thermodynamic equilibrium between the

intrinsic lattice defects. Through this influence

the ionic conductivity and the self diffusion coeffi- cients might be modified. It must, however, be kept

in mind that this assumption is arbitrary, and that

it is not justified in many practical cases. To explain the methods that can be used to deter- mine the components of the electrical conductivity

separately, it is necessary to give a short discussion

of the movement of the charge carriers under the

influence of an electrical field and a concentration

gradient.

The movement of ions in a crystalline solid can

only take place by the movement of lattice imperfections. As a consequence, the transport can be

described in two different ways.

In the first way, an average mobility is attributed

to the main ionic components of the crystal. Al- though the value of this mobility depends on the

fraction of those ions that are in a favourable

position to move (e.g., interstitial ions), and on the

mobility of these ions, the latter are not introduced

as such in the analysis. Therefore, it is unnecessary

to have a detailed knowledge of the actual model

valid for the substance under consideration. The

transport coefficients (mobility and diffusion coeffi- cient) are considered as empirical quantities, the

gradients occurring in the transport equations are those of electrical and thermodynamic potentials.

This method was used by C. Wagner in his original treatment of transport in growing tarnishing

layers [2]. In the second way of treatment, attention is focussed on the moving species themselves, the

actual lattice defects present. The real transport

properties of these defects and their concentrations

enter into the transport equations. Thus a detailed

model is required. From such a model the mobility

could, in principle, be calculated. The chemical

potentials that are introduced into the equations,

are the so-called virtual chemical potentials of

the defects [3]. In his review of Wagner's equations,

Kroger compares both ways of description [4].

We will start using the description in terms of defects. The quantity jk is the electrical current

density connected with the flow of particles of type

k having a charge ZkQ, where q is the absolute value

of the electronic charge. The concentration of these

particles is Uk, their mobility Vk and their diff^usion coefficient Dk- The partial current density can then

be written:

jk = — nk\zk\qvk grad V— DkZkq grad rik-, (1)

V denotes the electrical potential.

Here it is assumed, that no cross effects take

place, like those present in the general transport

equations of irreversible thermodynamics [5]. This means that, apart from the interaction through

the electrical field, the currents and gradients of one particle species do not influence the flow of

other species. Furthermore, only isothermal situa- tions are considered.

We introduce the electrochemical potential rj^-, which is connected to the chemical potential pi^

and the electrical potential by

Inhere

T/A- = iXk + ZkqV,

fJik = fJLko + kT In Hk.

Furthermore, we use the Einstein relation:

\zk\q

We obtain:

(3)

(4)

Here (Tk = nk\zk\qvk is the partial conductivity asso- ciated with the particle species k. If the particles considered are not defects, but

lattice ions themselves, eq. (4) is still valid; 7]k is in that case not the virtual electrochemical potential

of the defects, but the real electrochemical potential of the corresponding ions. Since, not considering the sign, both electrochemical potentials differ only by a constant, their gradients are equal.

3. Electronic Transport

Now, for the moment, we confine our attention to the electronic component je of the current only.

This is composed of a contribution by electrons

and holes:

2. The General Transport Equation xhus

je=jn+jp

cr„ , J Je= — grad T7« - ^ grad r)p.

(5)

(6)

Here iq„ and — rjp are equivalent to the quasi Fermi

potentials of electrons and holes, as used in semiconductor physics. Following Stieltjes [6] we write

this in the form

je = grad T7„ — — (grad T}n + grad Tjp).

(7)

150
Now we assume, that the thermodynamic equiUbrium between electrons and holes is not disturbed

by the current flow. This assumption excludes

effects of injection or extraction, Uke those occurring near p-n junctions in diodes and transistors. This

equilibrium situation has the consequence that:

grad T7„ + grad rjp = 0.

So the electronic current becomes:

(Tn + (Tp J

CTe , CTp , je = —^ grad Vn = — grad Vn^-— grad rjp.

(8)

(9)

4. Ionic Transport

We proceed now, by assuming that two ionic

species also contribute to the current transport. Let them be characterized by indices 1 and 2. Carrying out the same procedure again, we find for the total ionic current density =ji +^2:

ziq

grad Tji + — ( — grad tji —- grad 172

V

(10)

We assume now that, at each point, the thermodynamic equilibrium between the ionic species is

This relation, introduced into (10), gives the result:

undisturbed by the current flow. This thermodynamic equilibrium can be expressed by a reaction

equation like:

|22|MI^.I++ |2,|;^l^^|- (11)

The equilibrium condition requires the equality of

the sums of the chemical potentials on both sides

of this equation.

(12)

(In the indices we write MX instead of Mz^z^)

This equation is also valid for the gradients:

|zi| grad )U2 + |z2| grad /ai = grad ^t^/A. (13)

For small defect concentrations \j.mx is constant and

consequently:

grad ju.A/,Y=0. (14)

Furthermore, the gradients in the chemical potentials may be replaced by gradients in electro- chemical potentials since the electrical terms bal- ance out. Thus we obtain:

\z\ I grad 172 + |z2| grad rji = 0.

. . . o-i + 0-2

, 7, =;i+;2 = grad rji — grad 171= — — grad t}2.

z\q z-iq

(15)

(16)

This is the ionic equivalent of eq (9). Note, that we tacitly changed over from virtual chemical potentials of defects to chemical potentials of ions.

X\^^^ + \z2\h

(19)

5. Total Current

Now the total electrical current density ji, flowing

in a mixed conductor when gradients in the electrical potential and in chemical potential are present simultaneously follows from (9) and (16):

ji=je+ji^— grad r}n-— grad tji q Ziq

(17)

This can be written:

. (Te+(T

- grad 7?„- Y (~ g^^^ + ^^'^ '^')- (18)

If the thermodynamic equilibrium between ions

and electrons is locally undisturbed, the following

equilibrium reaction equations can be considered:

{M* is a neutral metal atom, X* one of the metalloid,

/i • is an electron-hole, e' an electron). We then have

the relations:

At2+ |Z2| f^P ~" M--^'' \Zl — [Xm. (20)

These can be written in terms of gradients, and

electrochemical potentials can be introduced. We

obtain:

grad tJLM= grad fit + |zi| grad (Xn = grad rji + \zi\ grad iq„, (21) grad ^tj- = grad 1x2+ I22I grad fXp = grad 172 + \z2\ grad 17,;.

151
Under the assumed conditions eq (18) reduces to

jt = ^ grad T7„ -— grad ^A/ = -— grad rjp-— grad fXx qzi q qz-z (22)

Here use is made of eq (15) and of the fact that zi

is positive and z-z is negative.

6. Electrode Processes

In order to be able to make use of the equations

obtained, it is necessary to consider a specific situation. In such a situation, the processes taking place

at the free surfaces or at the electrodes determine

the boundary conditions under which the current equations must be solved. Therefore we shall spend

a few moments to look at these processes in some

detail. From standard semiconductor practice we know

that transfer of electrons (or holes) from a contact

to a semiconductor is not always as easy as, for example, between two metals. We distinguish two

extremes. At so called "ohmic contacts," no ap- preciable potential drop arises when a current is passed in either direction. "Blocking or rectifying contacts" allow no significant current to pass in one direction. Intermediate cases often occur in practice. The behaviour of a contact can be de- scribed in terms of the work functions of the metal

and the semiconductor. In practice it is often diffi- cult to make contacts with desired characteristics.

With respect to the ionic currents, the same

problems arise as in wet electrochemistry. Apart

from energy barriers that may exist, equivalent

to those occurring in blocking electronic contacts

(activation polarization in wet electrochemistry),

here current hmiting may also take place as a con- sequence of the material transport accompanying

the current flow. When, for example, gas is evolved

in the electrode reaction, the difficulty in its trans- port away is even greater than in the case of a

liquid electrolyte, where the solubihty is generally

higher and where bubbles may be formed. The

equivalent of the latter would be destructive in the solid state situation [7]. Because of the low

solubility and diffusion coefficient of gases in solids, the electrode reactions are confined to the

immediate vicinity of the three-phase boundary:

electrode — electrolyte — gas phase.

The uncertainties in electrode properties do

not interfere seriously in the determination of the

total conductivity, since by the use of a four point

method or by measurements with alternating current, disturbing electrode effects can be avoided.

With d-c measurements, the electrode effects must be taken into account. However, these effects are not always a disadvantage. To the contrary

we will see that a proper choice of the electrodes can, in some cases, open the possibihty of separating

the electronic and ionic contributions to the conductivity. This was first recognized for the case

of silver sulphide by Wagner [8] and by Hebb [9]. The methods were summarized by Wagner [10, 11]. For the moment, we shall assume that, by the

choice of certain electrodes, we can determine the

electrode properties at will. We will apply the current eqs (9), (16), and (22) to four situations

where:

1. completely reversible electrodes are present

as well for electrons as for ions, so that the current

flow is not influenced in any respect, or

2. no current is supplied to the sample externally, or

3. only electrons are allowed to pass through the

sample via the electrodes, or

4. only ions are entering and leaving the sample

through reversible electrodes.

These different situations can often be approximated in practice often to such an extent, that meaningful interpretation of measurements in terms of the formulae to be derived below can be

made. In the following we will consider oxide

samples only.

6.1. Completely Reversible Electrodes

Samples with completely reversible electrodes

are represented schematically in the figures la

and lb. In figure la the sample is in equilibrium

with the electrodes consisting of the metal component M of the oxide MO; in figure lb it is in equilibrium with the oxygen present in the sur- rounding atmosphere through the porous inert electrodes. In the latter situation, the composition

of the oxide can be changed by variation of the

oxygen partial pressure of the atmosphere. Thus,

the influence of this partial pressure on the conductivity can be studied.

For evaluation of the current densities eqs (9) and

(16) are valid. Since the assumed reversibihty of the

electrodes has the consequence, that the oxide stays

in equilibrium everywhere in the sample, no concentration gradients are formed. This means, that grad 7} in eqs (9) and (16) can be replaced by zq grad

r (see eq (2)).

ji = — (Ti grad V, and je = — (Te grad V. (23)

This is simple Ohm's law behaviour.

152
M MO M

+

porous Pt

M"

\ +

Figure 1. Mixed ionic and electronic conductor between reversible electrodes.

a. The metal component M used as electrode; b. Porous inert electrodes in an atmosphere with defined oxygen partial pressure: c. Transport by metal ions causes weight changes

of anode and cathode tablets.

The ionic contribution to the current can be found

by determination of the weight i ncrease of the

cathode and the corresponding decrease in weight

of the anode in the situation of figure la or, in figure lb, by the measurement of the oxygen taken

up or evolved at the corresponding inert electrodes. In this way the ionic transport number is found:

O-f+CTe ji+je

(24)

Note that in this way no information is obtained

about the ionic specie which is responsible for the

transport (M" or O"). To get this information the

method of Tubandt [12] can be used. Here the sample is composed of two (or more) tablets pressed

together. By weighing before and after current flow

the transport of matter is determined (fig. Ic). When only oxygen is mobile, no weight changes

take place, mobile^etal ions cause an_increase in weight of the cathode tablet. Similar effects can be

seen if inert markers are used.

An important matter of consideration is to what

degree the ideal reversible electrode can be ap- proximated in practice. It will be clear, that there

must be a driving force to get a flow of oxygen into the sample at the cathode and out of it at the anode

(fig. lb). So, fundamentally, the boundary regions

cannot be in exact equilibrium with the gas phase.

This deviation from equihbrium depends strongly on the structure of the contact material and sample

surface. They determine the effective length of the

three phase boundary, where the electrode reactions

take place and the ease of diffusion of the gas to

this region. Slow gas diffusion causes stronger

deviations, the higher is the current density and the

lower is the gas pressure.

It must be recognized further that, for instance

at the cathode, O2 molecules must be spht up into atoms and then be ionized. Thus the electrode must

be an efficient catalyst for these reactions, whatever

the physical meaning of this statement may be.

Since these reactions are thermally activated processes, a strong temperature dependence of their

effect must be expected. This is in contrast with gas

diffusion limitations, which exhibit only a slight temperature dependence.

The flow hmitations mentioned cause polariza^v tion voltages to be built up in the electrode regions.

^THese can effect the apparent conductivity measured. Moreover, if reversibihty of the electron

transfer is assumed, such a polarization voltage

increases the electronic component of the current

flow at the contact. Since the effects will in general

not be equal at both electrodes, the current flow will cause a net loss or gain of oxygen. This change in composition influences the conductivity components again, so that erroneous results will be

obtained.

The occurrence of these phenomena can often

be detected by a gradual change in the current

after the application of a constant voltage or. in- versely, a gradual change in voltage after the

switching on of a constant current. An experimental example is given in figure 2. ^ Here the voltage between two thinly sputtered

Pt electrodes on a sample of stabilized zirconia

(15 mol percent CaO) is recorded against time.

Figure 2. Experimental stationary voltage vs. time (a) and

voltage vs. current {b) curves of a Zr02-CaO sample in air with contacts acting as reversible electrodes at low current density and showing polarization at higher current densities

The polarization causes oxygen loss and as a consequence additional electronic conduction (drop in a-c resistance).

^ Details will be published elsewhere.

290-621 0-68— 11 153
Ambient was air at both electrodes, the temperature

550 °C. A second curve gives the resistance which was measured simultaneously by means of a small

superimposed alternating current. Both curves indicate that a decrease in bulk

resistance takes place during current flow, which

must be attributed to a loss of oxygen. ^ This oxygen

deficiency results in the appearance of appreciable

electronic conductance. (At the oxygen pressure

used the initial conductivity was at least for 99.99

percent ionic.) This electronic conductivity was

confirmed by a determination of the amount of

oxygen transported. The latter showed a corre- sponding decrease in time, like the resistance.

This experiment shows that the speed of take up

of oxygen at the cathode may be limiting. The

magnitude of the effect depends strongly on the

electrode preparation, and it is only present at relatively high applied voltages (or current densities) and low temperatures. Moreover, it is more

pronounced at lower oxygen pressures (at the

cathode). The effect is also clearly evident from a current-voltage plot using the final values (fig. 2b).

The strong temperature dependence at a certain current density indicates that a thermally activated

process like 02-dissociation or soHd state diffusion

is the cause, rather than a limited gas diffusion rate.

6.2. No Current Supply. Sample Between

Different Oxygen Partial Pressures

We will consider an oxide sample which separates two regions where constant oxygen partial

pressures are maintained, and where the boundaries of the sample can be considered to be in equilibrium with the corresponding oxygen partial pressures. In figure 3a an example is sketched;

another example can be a layer of oxide on a metal

(fig. 3b). Electronic contacts may be present to measure the potential difference, but are not used

to supply an electric current. Under this condition, jt in eq (22) is zero, so that

also using eq (8) and putting Z2 = — 2 we obtain:

grad = grad fXo = -^ ti grad fXo

-J

grad /Xo2- (25)

Integrating between the boundaries I and II, we

can obtain a value for the difference in the electro- chemical potential of the electrons at both sides. This difference is just — q times the value of the

voltage that can be measured between two identical

Figure 3. Galvanic cells with an oxide as electrolyte acting between different oxygen chemical potentials

a. Differing oxygen pressures at both sides, b. Oxygen potential fixed by metal

electrode on one side. The latter is the case, when an oxide layer is growing on a metal.

metal electrodes (the chemical potentials and de- rived quantities are expressed per atom here in- stead of the usual gramatom). We thus obtain:

E = drj„=- —

Q Ji 4g

^djLt,02- (26)

If the oxygen potential difference is small, so that

ti can be considered as a constant (equal to an

average value U), we get simply:

or, since fXo^= /Xq^+ kT \n po2

4-q p,

(27)

(28)

^ See section 7.1.

These are the well known expressions given by

Wagner [2] for a galvanic cell with a mixed con- ductor as electrolyte. They allow a direct deter- mination of the ionic transport number from a measurement of the type represented in figure 3a. To determine the oxygen pressure dependence of the transport number, both pressures in the compartments I and II have to be varied simultaneously,

but, in order to use eqs (27) or (28), only a small

relative difference between pn and pi is allowed.

With a cell of the type of figure 3b this is not

possible.

The oxygen dependence of the transport number

can still be obtained with this set-up when £ as a function of pu is determined. This can be seen as

follows. Writing eq (26) in the form containing

154

2
AT In instead of ^jlo^ and differentiating with

respect to the upper integration Hmit we obtain:

Mln p«2

' kT

4?

(ti) at Po^=Pn- (29)

So the ionic transport number and its pressure

dependence can be obtained from the slope of an

experimental E versus In Pu plot.

Often the oxygen pressures in the two regions

I and II are fixed by fiUing them with a mixture of a metal with its oxide. This provides an easy ex- perimental technique in which the specimen is hghtly pressed between two different tablets composed of the two metal-metal oxide mixtures and

placed in an inert gas atmosphere in a furnace.

The tablets can serve as electrical contacts. The

oxygen chemical potential difference can be found

from the thermodynamic data of the oxide used

[15-21].

However, only a Umited number of oxygen potentials are available in this way, and the oxygen

pressures differ very often by many orders of magnitude. So the average transport number so obtained

from eq. (27) is a rather inexact quantity, unless it happens to be unity.

Moreover, if the transport number < 1 an ion current flows, and the oxide mixtures have to supply

or take up oxygen. Since the electrode reactions

take place at the interface, diffusion processes in the metal-metal oxide tablet must prevent ex- haustion of one of the components there. At moderate temperatures this is impossible and the

method fails [22]. A check can be obtained by building up a cell with the same electrode tablets, but with a purely

ionic conductor as an electrolyte (e.g., ZrOz — CaO).

A current is passed through this cell that gives a mass transport comparable with the expected value

in the open circuit situation of the original test

cell. Only if the emf. assumes a stationary value

immediately after the interruption of the current,

and independent of its direction of flow, meaningful

values of the transport number can be obtained

from the original cell. The author's experience

with these cells at temperatures below 1000 °C

have made him rather pessimistic about the value

of the method for quantitative measurements.

However, this method remains useful to ascertain

whether an ionic transport number is unity or not

at any pressure in the range between the equihbrium pressures of the tablets.

It may be worth recalling, that using an electrolyte in which it is well estabhshed that ti=l in the region of interest, the tablet method can be

used with advantage to determine thermodynamic

data of one of the electrode systems [15, 23-29].

The set-up of figure 3a requires an inert wall

forming a gas-tight fit to the sample in order to separate the two gas compartments [30]. This is also necessary to guarantee the one-dimensional

geometry that was assumed in the calculation.

Since this is difficult to realize in most cases,

the set-up of figure 4a has been chosen by various

workers. The sample is clamped between the ends

of ceramic tubes, through which the reference

gases can flow. Here the requirements for gas

tightness are not so high, since the reference oxygen

pressures are determined by streaming gas mixtures. However, effects at the outer boundary, where

the gas pressures are not well defined and where

surface properties may be decisive, may spoil the

homogeneous field distribution inside the sample.

Mitoff [31] avoided these complications by the use of a guard electrode on which a potential is imposed equal to the measured potential in the center of the sample (fig. 4b). The influence of

electronic currents through the gas phase, which

at high temperatures result from thermionic emission [32], is also minimized by this arrangement.

Another practical difficulty met in this kind of

experiment at high temperatures is the occurrence

of stray thermal emfs. These may be considerably

higher than the electrochemical potentials to be

measured. Therefore, they must be reduced as much as is possible by careful design of the apparatus and control of the gas flows. However, they

may not be completely eliminated in many cases, and means to change the reference atmospheres

from left to the right and visa-versa are required to separate the thermal and electrochemical effects.

It must be kept in mind that both the thermal emf

and the transport number depend on the oxygen

pressure. This means that a simple correction by

substraction of the voltage obtained with identical

atmospheres at both sides is not Justified.

electrode somple

i ^ -

V / /////// //r. V / /////////JJ

'.-^ inner contact

\ guord ring

Figure 4. Practical galvanic cell arrangements with the oxygen

pressures obtained by gas streams

a. Two-contact arrangement, b. Guard ring added to avoid surface effects.

155
A more fundamental difficulty lies in the assumption, made in the derivation of the equations, that the boundaries of the sample are in thermodynamic

equilibrium with the corresponding gas phases. The

fact that no current is supplied externally is not relevant as before, since, iff, # 1, a material trans- port still takes place through the sample: only the sum of the currents > and ji equals zero.

If for instance the surface reaction between oxy- gen in the gas phase and in the solid is slow, and

from experience with oxygen isotope exchange it is known that this is often the case [33-37] , jumps in the thermodynamic potential of oxygen may occur

at the surfaces. In such cases the measured emf.

will be too low. A way to check if this effect interferes would be

to repeat the same measurement with different

sample thicknesses so that the oxygen flow is varied

with the same chemical potential difference. If necessary, extrapolation to infinite thickness would

give the correct result. The value of the partial currents ji = —je flowing

through the sample follows directly from the eq (9) when the value for grad rje as found from eq (25)

is inserted

je = -ji = — grad ri„ = ti grad /x,,,. (30)

This can also be written:

o-/ , 1 aide

—Je =Ji = -7- teti grad )U,o, = —

45 4-q (Ti + (Te grad

(3i;

In the arrangement of figure 3b this current is pro- portional to the rate of growth of the oxide MO on

the metal M.^ Eqs (31) are forms of the well known

Wagner's equations [2] for the rate of growth of

oxide layers on metals and leads to a parabolic rate law since grad ^to2 at constant oxygen pressure

changes inversily proportional to the thickness of

the layer. From the equations it can be seen that the

lower of the two types of conductivities determines

the oxidation rate. So ionic transport numbers can

be determined from this reaction rate if the oxide

is a predominantly electronic conductor.

In the opposite case where te < ti the transport

of electrons is rate limiting and can be determined

from the material transport. In a setup like that of

figure 3a and measurement of the oxygen permeation the electronic transport number in stabilized

zirconia and hafnia can be determined.^ The avail-

^ Here, and in what follows, we assume that transport of material in uncharged form

does not take place.

^ See section 7.1.

ability of gas tight tubes of these materials makes

such a determination of the electronic conductivity

(through eq (30)) in them very easy: even the pres- ence of electrodes is not necessary.

6.3. Only Electrons Allowed to Pass Sample

6.3a. Voltage Chosen to Stop Ion Current An important variation of the last mentioned kind

of experiment consists of the application of a volt- age (or, what is equivalent, the passage of a current)

that is chosen in such a way that the growth of the

oxide layer or the permeation of the oxygen is just

stopped. In that situation, instead of the total cur- rent being zero, only the flow of the ionic component

is stopped. This means that according to eq (16):

grad r\ \ = grad 7)2 = 0. (32)

In combination with eq (21) and eq (8) this has the

consequence that

grad T7n = TT = :^r~\ grad fi02- (33)

The applied voltage, ^stop^ can thus be found as

(see eq (26)): -If" 1 2|z2|g

(34)

By comparison with eq (27) we see that this voltage is equal to the emf. that would be devel- oped, if the oxide were a purely ionic conductor.

Furthermore, since the current flowing is all elec- tronic, an average electronic conductivity can be

found directly (from eq (9) in integrated form).

An experiment of this kind is described by Jor- gensen [38] , who studied the oxidation of silicon. As has been stressed by Kroger [39], this kind of

experiment allows an unequivocal determination of

the cTiarge (|z2|) of the moving ionic specie. This

is important because it is not self evident that the

defects responsible for the current transport have

the charge which corresponds to the conventional

charge attributed to the ions.

It must be remarked, that, on behalf of clarity, this was tacitly assumed in many of the equations

used so far.

6.3b. Blocking Electrode(s) for Ionic Currents A second way to block the ionic current in oxides

is the use of contacts that are impermeable to and

do not take up or supply oxygen. The application

of a small voltage causes only a temporary ion flow.

156
After some time a concentration gradient is built up and no ion current flows any more: the cell is polarized. In this polarized condition the electron current is responsible for the complete electric current, just as in the situation of the last section.

So eqs (33) and (34) are also valid here, only the

potential difference Vn — V^ is the applied voltage

here and the difference in ^tj^ and [J-q^ is estabHshed automatically as a consequence of the ap- phed voltage and the damming up of the mobile

ions. The current in the final stationary state is given

by eq (9)

y,=ye = ^ grad T?„^ (35)

This can be written also, according to eq (33) and

since |z2| = 2:

jt =je = grad txo^. ( 36

)

Here CTp is in general dependent on fXQ,, cre = (Teip-o^), and thus dependent on the appUed voltage. Inte- gration over the sample thickness L gives:

1 f"

7(^ = — ^ j

CTp (yLloJ d/Ao^. (37)

Only for small applied voltages, corresponding to small differences in /Xqj. a meaningful average elec- tronic conductivity o-p can be put before the integral, resulting in:

^

(38)

However, one inconsistency is present in our treatment so far. By means of the applied voltage

only the difference /u,^ — /Mq^ is determined, and not

the value of each of them. Because of the perfect

blocking action of the electrodes, as was assumed

in the beginning of this section, the possibihty of

fixation of the /I02 by exchange with the surrounding

gas atmosphere is excluded. So the oxygen potential is only determined by the initial composition

of the sample, and depends on the way of its pretreatment.

To obtain meaningful results, one must be certain

that, during the establishment of the stationary,

polarized state, no oxygen loss or gain whatsoever

could have happened. However, this certainty is hard to obtain: in practice it is likely that some

change in composition occurs. Therefore it is ad- visable to use one blocking contact only, and allow

the sample to come into equilibrium with the atmos- phere by the use of a reversible electrode for the

other contact. Of course a metal-metal oxide mixture or the metal component of the oxide itself would do as well. We assume now that electrode I is chosen as the

reversible one and that, consequently /u,q^ has a well

defined value. Then ^t"^ is fixed by the apphed volt- age (see eq (34)):

^^o. = ^^o^^^E. (39)

In the same way, the oxygen potential at a position X in the sample, where, with a metallic probe,

a voltage E{x) is measured with respect to the

reference electrode I, is found as

l^o.= lA^ + ^qE{x). (40)

Using this equation in eq (36) and solving for (Xe()U.-^J gives

This probe method was introduced by Hebb [9] in the case of AgaS.

Wagner [10] suggested a method to obtain the

/i.02 dependence without the use of probes. Dif- ferentiation of eq (37) with respect to the upper

integration limit (introducing eq (39) and solving

for o-e(/i{]J gives

CTe{K,)-L^. (42)

So the electronic conductivity component is ob- tained from the slope of the j versus E plot which

is measured in the final polarized situation. The

corresponding {Xq^ is given again by eq (39). The integration of eq (37) can be performed, if

it is assumed that the electron (or hole) mobihty

is independent of the composition of the substance.

In that case, the oxygen pressure dependence of

CTe results only from the concentration dependences

of the electrons and holes [n and p). According to eq (2), n can be expressed in terms of )U,„:

ra(jc) = n, exp — (43)

Here ni is the electron concentration at the re- versible electrode I, and n{x) that at the positions.

With the eqs (33) and (40) this can be transformed

into:

, . -qE(x)

n{x) = ni exp —-pf (44)

157
Since the electron conductivity o"„ is proportional

to ra, we can also write: — qE{x)

(Tn{x) =(Tni exp (45)

For holes the equivalent expression is found:

qEjx)

ap(x) = (Tpi exp —j^— (40)

So, (Te{x) in eq (37) can be replaced by

-qE{x) qEjx)

(Te= Crn + (Tp= CTni CXp Jlj; h CTpi exp —

(47)

When also E is introduced for ixq^, we find, after integration of eq (37) between £ = 0 and E^E,

jt=je =

A typical shape of the curve representing this j versus E dependence for the polarized state is given in figure 5. It shows three regions: a low

voltage region where the electronic conductivity

is ra-type, a relatively flat portion corresponding

to intrinsic electronic conductivity, and a second

rising part at higher voltages corresponding to high

oxygen pressures and p-type conductivity. The

slope of the curve is a measure of the electronic

conductivity according to eq (42). It must be re- marked that in many practical cases only a part

of this general curve can be realized because of

the limited range of oxygen pressures in which the

E

Figure 5. Typical stationary current versus voltage curve for

a polarized cell The oxygen pressure at the reversible reference electrode is assumed to be so low

that the current is n-type when a low positive voltage E is applied to the blocking

electrode. At a higher voltage, part of the oxide becomes high ohmic (electronically). At still higher voltages, the region near the positive contact becomes p-type.

oxide under consideration is thermodynamically

stable. The polarized cell technique described has been

successfully applied to the silver, copper and

lead hahdes [40, 41, 42]. In applications to oxides

the determination method consisting of the passage

of a constant current and observation of the final stationary voltage has been described [43, 44, 45].

However, a reversible reference electrode was

not used in these cases. An analyses according to eq (48) was applied to Zr02.CaO and Th02.Y203

by Patterson et al. [46].

It seems useful to discuss the limitations of this method as applied to oxides in some detail. As has

already been discussed, the fixation of the oxygen

potential at one electrode is important. For instance,

it is unsatisfactory to make measurements in a vacuum, as was done by Danforth and Bodine [43] in a study of Th02-crystals. The oxygen partial

pressure of a practical vacuum is undefined.

An interfering process takes place at the free sides of the sample, on which no contact material can be present to close off the surrounding gas. The uniformity of the gradients inside the sample

may be spoiled and 02-exchange with the atmosphere takes place. Thus the current in the final stationary state may not depend only on the applied

voltage, but also on this oxygen exchange rate. In most practical cases it will not be possible to close

off the sides and the only measure that can be taken

to reduce this disturbance seems to be the use of

thin samples and, possibly, a guard ring technique.

Furthermore, the degree of blocking of the elec- trode is difficult to ascertain. Any leak of oxygen

through the contact gives rise to a stationary final current, just as that arising from the electronic

conductivity. The temperature dependence may,

sometimes, give an indication of what process

might be responsible, as shown by Vest and Tallan

[45], but cannot be used as a general criterion. It

is evident that the kind of electrode and its way of

application are very important in this respect.

This is demonstrated in figure 6. Shown are some

recordings of the voltage versus time when a constant current (10 ^lA/cm^) is switched on and off with three different kinds of Pt — electrodes and

with different oxygen pressures. The sample was

calcia-stabilized zirconia, its conductivity at the

temperature concerned (550 °C) was less than

10~^ percent electronic (determined by the oxygen

permeation technique).

Note the difference from the experiment shown

in figure 2 — the coordinate values are in millivolts and seconds; in figure 2, in volts and minutes. The

contacts used in the former experiment behaved

like those in figure 6a with small currents and thus

158
0.01 0.1 1 10 IOC torr ^ ^

n.

0 ^0 l.mKsec ) »

Figure 6. Experimental current vs. time curves for ZrO 2-CaO

during a current pulse of40 seconds and 10 pAlcm"^

Different types of electrodes and different oxygen pressures were used. The building up of a polarization voltage occurs only when some ion blocking takes place: a. re- versible electrodes of thinly sputtered platinum; b. more densely sputtered platinum

electrodes block at high and low oxygen pressures; c. painted platinum electrodes show more blocking action.

were perfectly reversible in this current and volt- age range.

In conclusion, we can remark that the polarization technique, which is in principle capable of

the detection of very small electronic contributions to the conductivity, seems in the practical

apphcation to oxides only dependable for the de- termination of substantial electronic transport

numbers. Results can only be trusted quantitatively if care has been taken to fix the composition

of the sample, if only low polarization voltages have

been used, and if the use of different blocking

electrodes gave consistent results.

6.4. Only Ion Current Allowed

While, in the preceding section, ion current blocking was discussed, now the blocking of elec- trons will be considered. This can be achieved by

the use of a purely ionic conductor, exhibiting negligible electronic conductivity, as an electrode. The

situation is most simple if the ionic conductance of

this electrode is much higher than that of the sample

under test. Then, when a current flows, the voltage

drop over this electrode caused by the current can be neglected. This intermediate electrode, as well as the opposite side of the sample, are pro- vided with reversible electrodes. These are in contact with the same surrounding, oxygen-containing,

atmosphere (see fig. 7a). If the applied potential difference Vu—Vi^E is so small that no electrolysis products can be formed

at the interface, the distribution of two different potentials is as represented in figure 7b. Since in the sample je — 0 it follows from eq (9) that At);, = 0. So the total applied rjn-difference {= — qE) appears

a

b

Figure 7. a. Blocking ofelectronic current by an ionic conductor

as electrode.

b. Distribution of oxygen chemical potential and electron electrochemical potential

in sample (right) and ionic electrode (left).

over the ionic conductor. Consequently, a drop in oxygen chemical potential of A)u,02~4g£' builds

up over it. As both contacts are in the same atmosphere, a corresponding inverse oxygen potential

difference appears over the sample.

The ionic current through the sample is given by:

ji = ^ grad 172 (see eq (16)).

Furthermore, according to eq (21):

grad 172 = i grad (Mq^ — 2 grad i7„.

Because here grad 17,, = 0, this becomes grad

172 = 1 grad ^02- Thus, for the current density we

find:

ii^^ grad /xoj. (49)

If cTi is independent of the oxygen pressure, or if the applied voltage is sufficiently low so that cr, can be considered as constant, integration gives

So, it is just as if no electronic conductivity existed

in the sample, and as if the voltage was applied

directly. In fact the ionic conductor has the function

of presenting an oxygen atmosphere (different from

that of the surroundings) and the electrical current

is a measure for the oxygen penetration through

the sample which occurs as a consequence of the

applied oxygen pressure difference.

159
Practical difficulties are likely to be met again by

the processes taking place at the sides of the sample

and with exchange of oxygen between the sohdsohd interface and the surrounding atmosphere.

They have to be solved by closing off the sides or, perhaps, by the use of a guarded electrode arrange- ment. So far no studies are known by the author

where this method has been used in oxide research.

An important variation results when electrode

II is blocking for ions. Any current passed through

the system enriches or impoverishes the sample

with oxygen. In this way, by measuring the equilibrium voltage (= equilibrium pressure) after a certain

amount of charge has been passed as a function

of this amount, the solubility of oxygen as a function

of the oxygen pressure can be measured. This

method of "potentiometric titration" has for instance been used by Alcock and Belford to deter- mine the solubihty of oxygen in liquid lead, and

recently, by Alcock, Zador and Steele to study the

nonstoichiometric oxides TiO-z , NbO-i, and Ti Nb; O2

[47].

7. Survey of Existing Data

At present, dependable quantitative knowledge

about ionic conductivity is only available for a hmited number of oxides. Only a short review will be given here and it will be restricted to well

estabUshed data for some important refractory

oxides.

7.1. Stabilized Zirconia

The literature on stabihzed zirconia is very

extensive and still rapidly growing. The ionic

conductivity of rare earth doped zirconia was al- ready recognized by Nernst in the year 1899 [48]. Schottky in 1935 [49] suggested the use in fuel cells and Baur and Preis [50] investigated this application experimentally in 1937. Their emf measure- ments proved the predominant ionic character of the conductivity. Wagner explained this by the

assumption of the presence of mobile oxygen

vacancies (1943) [51]. Confirmation was obtained

by the comparison of x-ray and pycnometer densities by Hund [52] and by comparison of oxygen

self diffusion data with the conductivity [37, 53].

The correlation factor [37] appeared to correspond

to a vacancy mechanism. The cation mobility is very low [54]. The influence of the dope concentration on the

conductivity has been studied extensively [55-67].

It appears that a maximum in the conductivity occurs somewhat above the lowest doping level that is required to make the cubic structure stable. For calcia the maximum is situated at 12 mol per- cent CaO. For rare earth doping, it hes between

8 and about 12 mol percent. The conductivity maximum shifts to higher percentages in the direction

Sc, Yb, Gd, Sm, Nd, i.e., in the direction of higher

ionic radius and with higher temperatures. The

highest conductivity is found for the smaller

ions (Sc).

The activation energy is minimum for the highest

conductivity and increases from there with higher

concentrations and also with increasing ionic

radii [57, 58, 62].

From this experience it is clear that the ionic

conductivity is more determined by the mobility

of the oxygen ions, than by the vacancy concentration. At 1000 °C slow irreversible changes take place

in the lattice which are accompanied by a decrease

in conductivity. Annealing at about 1300 °C re- stores the original conductivity, which has remained

ionic in character [63-67]. X-ray and neutron diffrac- tion studies indicate that some ordering mechanism

takes place. A practical study by Takahashi et al. [66] reveals that for an Yb203 or Y2O3 content above

10 percent, anneahng at 1000 °C does not influence

the 1000 °C conductivity. But the 600 °C conductivity decreases somewhat after a 1000 °C anneal.

The ionic transport number measured by direct

transport and emf. methods appears to be 100 per- cent within experimental error at near atmospheric

pressures, and also at such low oxygen pressures

as coexistent with H2O+O2 or CO+CO2 mixtures or with metals hke Fe, Ni, Cr [15, 68, 69]. But at very low pressures, e.g., as low as in equilibrium with metallic Zr, the transport number

drops because the increased electron conductivity

predominates [21, 30, 46, 70, 71, 72]. This effect can be observed directly by transport

measurement of zirconia in a reducing atmosphere

[14] and is most easily demonstrated by trying to electrohze oxygen out of a gas system which contains no oxygen or reducible compounds. The

zirconia is reduced itself in this way, but, at low

current densities, no metalhc Zr is formed. This can be seen from the equilibrium potential after the current is interrupted, which is lower than

would correspond with free Zr [72]. The drop in resistance shown in figure 2 is caused by the same

effect. The resistance increase found by Strickler

et al. [61] upon electrolysis in a reducing atmosphere

(Zr02-Y203) is not understandable in this respect,

and is opposite to the author's experience (with

Zr02-CaO).

The small contribution of electronic conductivity

at higher pressures can be obtained from oxygen

permeabihty measurements as were performed by

Smith et al. [73] as a function of pressure (dif- ference). The p^^'* law found can be explained from

the most simple defect model with hole conduction

160
and a concentration of ionized oxygen vacancies

fixed by the calcium concentration.

Analysis of Smith's data with the help of eq. (30) and resistance data from ref. [55] yields a transport

number for holes at 1275 °C and 25 torr oxygen pres- sure of about .05 percent. The activation energy

is 2.5 eV, so is about two times higher than that

for ion transport. Our measurements at lower

temperatures confirm the activation energy, but

yielded higher values. At 730 °C the electronic

transport number found was 3.10"^ for ZrOz with

15 mol percent CaO.

Predominant electronic (n-type) conductivity

appears with the addition of comparatively small

amounts of V2O5 [45] or Ta205 [74], which act as donors. A solution for the problem of why such a small donor concentration in the presence of a much

higher acceptor (Ca) concentration can give n-type

conductivity was proposed by Kroger [75]. He as- sumed that the greater part of the effectively negatively charged calcium ions form neutral associates

with the oxygen ion vacancies.

In this way, the acceptor character of the calcium

is annihilated and only a small concentration of

five-valent ions acting as donors are sufficient to compensate the rest of the (unassociated) calcium

ions and give rise to «-type conductivity.

The oxygen movement takes place through the

vacancies neighbouring calcium ions and is possible because of the high concentrations present.

A Mossbauer study [76] of iron-57, and mechanical or electrical loss measurements [77] (for Th02-

CaO) point to the presence of acceptor-vacancy

associates.

7.2. Pure Monoclinic Zirconia

The low temperature modification of zirconia is monoclinic. Its electrical conductivity was studied

by Rudolph [78], by Kofstad and Ruzicka [79] and,

most recently, by Vest, Tallan and Tripp [80]. According to the last mentioned investigators the

conductivity is mainly electronic and is proportional

to Pq'^^ at pressures above 10~^ atm. The conductivity curve shows a flat portion between 10"** and

10"^' atm. and increases towards lower pressures. In the defect model proposed, completely ionized

zirconium vacancies are assumed. However, the

possibility of the existence of such highly charged

defects was doubted by Kroger [75], who argues

that association with ionized oxygen vacancies is more probable.

7.3. Pure Tetragonal Zirconia

This modification is stable at high temperatures.

Electrical conductivity measurements performed

by Vest and Tallan [81] showed only a very weak

oxygen pressure dependence in the temperature

range 1300 to 1700 °C. A shallow minimum in the

conductivity-pressure curve was found near 10~^ atm. at 1300 °C, which shifts to higher pressures at increasing temperature. Polarization measurements

revealed mixed ionic and electronic conductivity.

The ionic contribution was dominant at 1300 °C and

practically pressure independent, the electronic

part showed a pronounced minimum, with n- and

p-type branches at low and at high pressures. Oxygen dependences in these branches approximated Pq^^^ and p^^ laws. At higher temperatures,

the electronic transport number increases to about

h (1600 °C).

7.4. Stabilized Hafnium Oxide

The properties of calcia stabilized hafnia were

investigated by Johansen and Cleary [57], yttria

stabilized hafnia was studied by Robert, Deportes

and Besson [82]. As is to be expected, the con- ductive properties are very similar to those of

stabilized zirconia. Even the slow irreversible

changes near 1000 °C take place in Y-stabilized

hafnia [82]. From the oxygen permeation data of Smith

et al. [73] it can be calculated that the electronic

t ransporj; number is also of the s a me order_Qfjnag;

nitude as in zirconia. Pure monoclinic hafnia can be studied in a greater

temperature range than monoclinic zirconia be- cause of the higher transformation temperature.

Tallan, Tripp and Vest [83] found, as with zirconia,

electronic (p-type) conductivity at oxygen pressures

higher than 10-« atm. (between 1000 and 1500 °C).

The pressure dependence was again as p^. The

ionic contribution was low, as was checked by

polarization measurements with blocking electrodes.

7.5. Thorium Oxide

Thorium oxide exhibits mixed conduction, as was demonstrated by Danforth and Bodine [43, 44]

by a polarization technique. More quantitative

results were obtained by Rudolph [78]. By means

of galvanic cell measurements Subbarao, Sutter and Hrizo [84] found that the ionic transference

number of Y-doped thoria at near atmospheric

oxygen pressure was lower than unity and dependent on the Y-content. Alcock and Steele [21, 30]

established a value of unity at low oxygen pressures. Their galvanic cell measurements demonstrate that this is also the case at such extremely

low oxygen partial pressures that stabilized zirconia

would become electronically conductive.

Results by Lasker and Rapp [22] and Bauerle

[85] show that a pressure independent conductivity

plateau exists at pressures below about 10"'' atm.

161
.of oxygen. This corresponds to purely ionic conductivity. At higher pressures p-type conductivity comes up, increasing with a V4 power of the pressure. The ionic conductivity plateau increases

with dope (Y) concentration up to about 15 mol/YOi.5

and decreases with higher dopant content. The

resuhs for "pure" thoria indicate that the properties

are still determined by the presence of unknown

impurities.

Density measurements performed by Wimmer,

Bidwell and Tallan [86] indicate that oxygen va- cancies are responsible for the ionic conductivity.

The activation energies for the electronic conductivity and for the ionic conductivity are both about

equal to 1 eV, so that the transport number does

not change appreciably with the temperature.

7.6. Cerium Dioxide

In the older literature about cerium dioxide

ionic conduction was sometimes suggested [87], but measurements were also interpreted in terms

of electronic conduction [88, 89]. Ionic conductivity was proved by Neuimin and Pal'guev [89] and

identified by mass transport measurements as oxygen conductivity [90]. The passage of a direct current may result in a reduction of the Ce02,

while at the same time the conductivity increases,

as was shown by Kevane et al..[91]. Cerium oxide

doped with La2 03 was studied by Takahashi, Ito and Iwahara [92]. A maximum in conductivity was

found near a content of 10 mol percent. Emf. meas- urements show that the conductivity is 100 percent

ionic at near atmospheric oxygen pressures if^ more

than 2.5 percent of La203 is present. At lower

oxygen pressures reduction takes place giving rise

to n-type conductivity.

7.7. Ternary Mixed Crystals

Mixed crystals in the system Zr02 — Ce02 — Y2O3

were studied by Robert et al. [82]. They found that the cubic structure was stabilized already with

3 mol percent Y2O3 when part of the Zr02 was re- placed by Ce02 (in the ratio Zr02 : Ce02 : Y2O3 = .90 : .07 : .03). However, the maximum conductivity still appeared at 8 mol percent, just as

in the case with zirconia. The value of the maximum

did not differ appreciably from the zirconia case.

7.8. Aluminum Oxide

The conductivity studies of alumina do not give

sufficient information about the ionic contribution

[100, 101]. Preliminary galvanic cell measurements

by R. W. Cooper on single crystals [102] indicate an appreciable ionic component. This is also found

by Schmalzried [71] for sintered alumina.

7.9. Yttrium Oxide

Cell measurements by Tare and Schmalzried

[103] at 825 °C indicate purely ionic conduction in the oxygen pressure range 10~^ to 10~'^ atm. However, measurements by the polarization technique

performed by Tallan and Vest [104] on purer

samples and at higher temperatures indicate only

a very small ionic contribution. This was even so

at the oxygen pressure where the conductivity is minimum.

7.10. Beryllium Oxide

The electrical conductivity and self diffusion

coefficients of BeO have been studied intensively

[18, 105-116]. It is well established now that the

conductivity is of mainly ionic character over a wide range of oxygen pressures and at temperatures up to_ 1600 °C [18, 110, 111, 112, 113], al- though mainly electronic conduction was proposed

by Pryor [109]. Comparison with self diffusion data

of oxygen and beryllium [105] indicates that the

beryllium ions must be responsible for the current

transport. The value of the correlation factor points

to a vacancy mechanism [111, 112].

The conductivity versus l/T curve shows an

activation energy of about 2.6 eV. It was assumed

that this conductivity is extrinsic and caused by

unknown impurities [112, 116].

At temperatures below 1000 °C, a much smaller

slope was found [112] and it was suggested that

in this region the current passed through a sec- ond phase formed at the grain boundaries. This

would result from impurities present in a concentration above the solubihty limit at the lower temperatures.

7.11. Magnesium Oxide

After some interpretation difficulties of earlier results [93], Schmalzried [94] and Mitofif [95] agreed

about the mixed character of the electrical conductivity of MgO. The ionic transport number de- pends on the temperature and on the oxygen

partial pressure [18, 96, 31, 71]. It is maximum at about 10 atm. oxygen and decreases when the

temperature is increased from 1000 to 1500 °C. Comparison of self diffusion data for oxygen [97] and for Mg [98] leads to the conclusion that the

total electrical conductivity is influenced by the

addition of Ni [99].

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164
Measurement of High-Temperature Thermodynamic Properties of NonStoichiometric Oxides Using Solid State EMF and Coulometric

Techniques

B. C. H. Steele

Imperial College, London, England

The relative partial molar free energy of oxygen in a number of non-stoichiometric oxides, Ti02,

VO2, Nb02, M0O2, UO2, has been measured as a function of composition and temperature (700 to 1100 °C) using galvanic cells incorporating solid oxide electrolytes. The data have been analysed to find the region of nonstoichiometry over vv'hich the dilute solution laws can be applied and the associated values for the partial heats of solution of oxygen are briefly reviewed. A coulometric technique was

successfully employed to vary the oxygen content of certain of the oxides, and the advantages and

present limitations to this method of controlling the composition of non-stoichiometric oxides are discussed with reference to the different cell assembUes and oxide systems examined. Key Words: Coulometric titration, free energy, galvanic cell, nonstoichiometry, solid oxygen

electrolyte.

1. Introduction

Thermodynamic studies of non-stoichiometric

oxides can help to clarify current theoretical

problems associated with the properties of these

phases of variable composition. In any metal oxide

system for example, it is important to know the

composition range over which the simple WagnerSchottky model for nonstoichiometry is applicable.

This model requires random non-interacting defects

(Henrian dilute solution behaviour) and is characterised by a constant partial molar heat of solution over the relevant composition range. If the model

is applicable then the postulated defect equihbria can also be expressed in terms of simplified massaction constants in which activities are replaced

by concentrations of the appropriate species [1].' These regions of complete disorder are, however,

difficult to establish experimentally as the partial

molar free energy-composition isotherms become

very steep when the metal/oxygen ratio approaches

the stoichiometric value (figs, la and lb). As the deviations from exact stoichiometry in- crease interaction between the defects can result

in the formation of 'micro-domains' in which the

defects may be considered to be ordered [2]. This

ordering process can produce a series of ordered

intermediate phases of definite compositions and

structures as exemplified by the Magneh struc- tures in Ti„02„-i, V;,02«-i. Under these conditions

the partial molar heat of solution is no longer an

invariant function of composition, and activity

' Figures in brackets indicate the literature references at the end of this paper.

coefficients should be introduced into the relevant mass action expressions, although many theoretical models have been proposed which do not comply with this requirement.

It is obviously important to obtain data on the

thermodynamic stability and composition range of

these ordered intermediate phases as their existence

is obviously fundamental to an understanding of mass transport processes in oxide crystals, and

thus to the mechanisms involved in ionic diffusion,

ionic conductivity, and solid state reactions. Thus

ordered intermediate phases, which often appear

only after long annealing times, seem to be only

slightly more stable than the disordered parent

structure from which they are derived. Moreover,

differences in the thermodynamic stability of the

individual intermediate phases also appear to be

very small [3, 4] and difficult to determine by conventional gas equilibration techniques using

CO-CO2 or H2-H2O mixtures.

In view of the preceding comments it is apparent

that use of a solid state coulometric technique to control the composition of a non-stoichiometric

phase offers many potential advantages. This tech- nique was employed by Wagner et al. to control

the stoichiometry of sulphides [5, 6] and tellurides

[7], and more recently has been used to study the

thermodynamics of UO. + x [8], TiOo-x [9J, NbOa-x

[9], FeOi + j. [10]. A titration, for example, of 0.01 coulombs (lOfxA for 16.7 mins.) involves the mass

transfer of 10"^ equivalents (8x10"'' gms) of

oxygen. For a 100 mg sample of an oxide MO pos- sessing an atomic weight of 100, this transfer of

165
T

-50 —

-100

-110 -

-120 -

I I I \ L_

0 0 01 0 02 0 03 0 04 0 05

COMPOSITION X in UO,. „

002 004 OOe 008 COMPOSITION X in MO,.,

Figure 1. Relative partial molar free energy (AGoP composition

isotherms for Ti02-x ond UO2 + Xo.t 1000 °C.

oxygen represents a change in the oxygen/metal

ratio of approximately 0.0001. Furthermore the emf

of these sohd state galvanic cells can often be

measured with a precision of ± ImV or better which

corresponds to an uncertainty of ± 100 cal. in the

derived partial molar free energy data for the

solution of oxygen.

700 800 900 1000 1100 1200 Temperature °C Figure 2. Oxygen chemical potentials of selected metal-metal

oxide systems as a function of temperature.

It is also possible using this technique to measure

oxygen partial pressures outside the range that can be conveniently established by conventional

gas equilibration techniques (see fig. 2). The low

oxygen chemical potentials associated with the

Nb-NbO [11] and Mn-MnO [12] systems depicted

in figure 2, for example, have been studied using

thoria based solid oxide electrolytes; the relevant

oxygen partial pressures at 1000 °C possessing

values of lO"^'*^ and lO^'^'*-^ atm. respectively.

To date, however, it has proved impossible to exploit all the potential advantages of this solid state coulometric technique and it is opportune

therefore to consider recent applications of this technique to non-stoichiometric studies in order to examine some of the current experimental problems.

2. Experimental

The procedures generally employed for high

temperature equilibrium emf measurements using

galvanic cells incorporating solid oxide electrolytes

have recently been reviewed elsewhere [13] and

it is only necessary therefore to describe those

additional features particularly relevant to the solid state titration techniques.

2.1. Selection of Electrolyte Material

Solid oxide solutions based on Zr02 or Th02

have invariably been selected as the electrolyte

material, and available data listed below clearly

indicate that in these materials the cationic contribution to the ionic conductivity is many orders

of magnitude smaller than the oxygen ion conductivity and may be neglected for most applications.

166
Zro.85 Cao.15 O1.85:

D,,.- (1000 °C)= 8X 10-8 cm^sec.-' [37]

D^r- (1000 °C) = 5 X 10-18 sec. [38] Dca- (1000 °C) = 3 X 10-18 sec. [38]

ThO-,:

D„- (1000 °C)

D. (1000 °C)

2 X 10-" cm 2 sec. -1 [39] 2 X 10-16 cm- sec. -1 [40]

The magnitude of the electronic conductivity

is more uncertain and is of course dependent

upon the oxygen partial pressure. Using the HebbWagner polarisation technique Rapp [14] has re- cently determined the electronic conductivity in

Zro.85 Cao.i5 Oi.85 and Thoss Y0.15 O1.923, and the

results for 1000 °C are summarised in figure 3. Oxygen permeabihty measurements of Roberts

and Fabre also indicate comparable electronic

transference numbers (10 ---10-^) at the higher

oxygen partial pressures. Particular attention

should be paid to these permeability measurements

which have included studies on selected commercially available electrolyte materials now being

used in many industrial and University laboratories; and it must be recognised that the flux of oxygen through these commercial materials is often much higher than would be predicted from

the electronic transference numbers determined

for similar electrolyte compositions prepared in the laboratory [15, 16]. Polarisation studies made at Imperial College

[17] suggest that the onset of significant n-type

conduction (> 1 percent) occurs at oxygen partial

pressures significantly greater than those re- ported by Rapp and a tentative value for this n-conduction in Zro.85-Cao.15O1. 85 has also been

included in figure 3. Accurate data for the electronic transference

number is not only necessary for precise coulometric

titrations but is also important because the pres- ence of an electronic conductivity allows a continuous flux of oxygen ions to pass through any

galvanic cell possessing a finite emf. For example

a Zro.85Cafl,i.50i.85 pellet of 1 sq. cm. cross-section

and 1 mm thickness will possess a resistance of

2.5 ohms at 1000 °C. If the observed emf is 500mV

then an electronic transference number of 10-- could

allow a current of 2mA to flow assuming that the

electrochemical reactions at the electrodes are not rate-controlhng. In the absence of appropriate

precautions this oxygen flux could significantly

change the composition of the non-stoichiometric

phase under investigation.

2.2. Reference and Working Electrodes

For coulometric titrations it is generally desirable to have separate reference and working

electrodes to avoid excessive polarisation at both

the working electrode and the relevant variable

composition phase electrode. It should be noted

that large concentration polarisation voltages may

cause reduction of the sohd oxide electrolyte thus

introducing an unknown electronic contribution

to the titration current.- To minimise polarisation

effects the electrode materials should exhibit both

large exchange currents and large diffusion coefficients to avoid concentration gradients. In- formation relating to the electrochemical kinetics

of systems involving solid oxide electrolytes is meagre but activation polarisation is probably not important, and attention should be devoted to minimising the concentration polarisation effects. Porous platinum oxygen electrodes, for example,

exhibit insignificant overpotentials [18] at 1000 °C even at current densities approaching lOOjitA cm

^ Using a pure oxygen reference electrode the voltage which can be applied across

a cell incorporating Zro.NsCaii.i.-.Oi.Hr, as the electrolyte should not exceed about 1 volt

at 1000 °C and 1.5 volts at 500 "C. Higher voltages will electrolyse the solid and give

rise to electronic conduction. By combining the oxygen chemical potential at which

significant electronic conduction (0.1%) occurs with that imposed by the reference

electrode, the maximum applied voltage can be determined using the well known

relationship; A/ii)j = — zEF.

Log c

-6 —

-8

-10

-30-6

{Ref17)

-10 -15 -20 -25

Figure 3. Ionic and electronic conductivities as a function ofoxygen partial pressure for Zrg g^Cao.isO , ss at 1000 °C [14].

167
The oxygen chemical potentials in equihbrium

with many non-stoichiometric phases of interest,

however, possess relatively low values, and as it is advantageous to avoid large chemical potential

gradients then alternative reference and working

electrodes must be used. Porous platinum elec- trodes in contact with CO-CO2 mixtures could be

used but present design problems if separate

working and reference electrodes are to be employed. Saturated liquid metal electrodes (e.g., Pb-PbO [19] or Sn-SnO [20]) are attractive be- cause the relatively high oxygen diffusion coefficients in liquid metals (~ 10"-^ cm- sec) ensures

small concentration overpotentials, and these elec- trodes can be incorporated into relatively simple

cell assemblies [21]. It is easiest, however, to specify solid reference and working electrodes,

and both Ni-NiO and Fe-FeO compacts have been

employed for this purpose. Cation diffusion rates

in the former system, however, are comparatively slow [22] (10-'"-10-i' cm^ sec'^ at 1000 °C)

and Ni-NiO electrodes can exhibit significant

overpotentials [23] at 1000 °C even at current

densities as low as 5/LtA cm"'-. Iron-wiistite elec- trodes (Dpe values in FeO vary from 10~^-10~^ cm- sec"' at 1000 °C) can sustain much higher

currents (~ 100/xA cm"- at 1000 °C) without excessive departures from the equilibrium potential

and are generally to be preferred, although re- action with the electrolyte material may sometimes

produce erroneous voltages. If the non-stoichiometric phase under examination constitutes the

other electrode then polarisation of this electrode

will often determine the maximum titration cur- rent which can be applied. For example during

studies involving coulometric titrations from an Fe-FeO electrode into a wiistite electrode of variable composition, currents as high as lOOO/xA

cm"'- were employed by Rizzo et al. [10] without

excessive polarisation effects. Coulometric titra- tion experiments by Alcock et al. [9] also involving an Fe—FeO working electrode but with non-stoichiometric rutile (D,,.- at 1000 °C: 10"'' cm'- sec"'

[25]) as the other electrode material, were limited,

however, to currents of 1-5/xA cm"-, and even

with these small currents subsequent equilibration

took at least 24 hours. This need to restrict titra- tion currents to small values for certain electrode

systems has lead to developments in the cell design

which are described in section 2.3. Another factor which should always be con- sidered when coulometric titrations involve solid electrodes is the influence of the three phase region

where the electrolyte, electrode, and gas phase

are all in contact. It is always possible that oxygen

ions discharged in the vicinity of this region could

form gaseous molecules which would not contribute to changes in the composition of the nonstoichiometric electrode phase.

2.3. Cell Assemblies

Typical cell arrangements that have been used

to examine the thermodynamics of variable composition phases are depicted in figures 4-6. Cell

assemblies I and II (fig. 4) have been successfully

used for many years at Imperial College for equilibrium emf measurements [11, 12, 26]. A special

feature of this design is the provision of two separate electrode compartments to eliminate gas phase transfer between the two electrodes via

CO and Ho impurities. The side reactions associated

with this process can produce erroneous mixed

potentials. Both designs (I and II) can be operated

with two separate purified streams of argon or under vacuum conditions (10 "'^-10"^ Torr).

Type I was adopted by U.K.A.E.A., Harwell,

for solid state coulometric studies on UO^ + j- [8] and the accompanying research programme was

successfully completed without any major practical difficulties. Studies on Ti02-j and Nb02-x

[9] using the type II cell configuration, however,

produced many experimental problems which were

principally caused by the excessive times required

for equilibration following a coulometric titration. Cell assemblies III and IV (fig. 5) were therefore

developed in attempts to overcome the limitations

I 12-15mm envelope surrounded by -earthed Pt snield

1. Non-stoichiometric electrode

2. Working electrode

3. Reference electrode

1 . Non-stoichiometric electrode

2 working electrode

3 . F?eference electrode WWWX Zrg COq ,5 electrolyte

///// Recrystollised ' Purox' AljOg

Figure 4. Cell assembly for solid state coulometric titrations (Types I and II ).

168
12 - ISmm

Metal - ceramic

seal

springs-"

1. Reference electrode

2. Pt - ©2 electrode www Alj O3

X^'^^^^ThQ.gYQ^ 0^.45 electrolyte

Ionic liquid (e.g sodium silicate.)

, - Variable composition phase

12-15mm

Electromicroba lance

sprmgs

1. Reference electrode

2. Pt - O2 electrode

3. Non-stoichiometric phase

^/////ZrQ.ggCaQ.^gO^as

Figure 5. Cell assembly for solid state coulometric titrations (Types III and IV).

imposed by the relatively slow solid state diffusion processes.

In the type III cell arrangement powdered

material consisting of the relevant non-stoichiometric phase was placed in an appropriate inert container and suspended from an electro-microbalance having a sensitivity of ±10~^ gm. The

compositions of both the powdered material and

a known volume of a recirculating mixture of CO

and CO2 mutually adjusted to their appropriate

equilibrium values and the associated oxygen

chemical potential was determined using the solid oxide electrolyte galvanic cell. By means of a coulometric titration a known amount of oxygen was then introduced into the compartment marked

/8 containing the non-stoichiometric phase. The

additional oxygen was distributed between the

non-stoichiometric phase and the CO-CO2 mixture.

The weight change of the powdered material was

recorded by the electro-microbalance, and the

oxygen chemical potential developed by the new CO-CO2 ratio obtained from the relevant e.m.f. measurement. All the measurements should be

consistent with the appropriate oxygen mass bal- ance values. It has been impossible, however, to obtain consistent reproducible data with this type

III cell arrangement due to problems associated

with the oxygen permeability of the zirconia electrolyte tube, and difficulties in selecting the optimum

quantities of non-stoichiometric powder and sur- rounding CO-CO9 mixture.

In another cell arrangement, type IV depicted

in figure 5, an attempt has been made to provide

the powdered non-stoichiometric phase with its own gaseous environment in an impermeable alumina crucible, and to alter the composition of this environment (CO-CO2, or H2—H2O) by coulometric

titrations using a porous Pt oxygen electrode,

which was connected to the external circuit via a platinum-alumina seal incorporated into the

alumina container. The relevant surfaces of the

alumina container and oxide electrolyte pellets were prepared with a final poHsh of V4 micron

diamond paste, pressed firmly together, and a

relatively impermeable seal was obtained by immersing the alumina-electrolyte interface in an

ionic liquid. A sodium silicate having the composition, Na20, Si02, 25:75 mole percent and melting

point ~ 800 °C was usually employed as the seaUng

fluid. Applications of this modified technique have

so far been confined to the Ti02-x system, and

preliminary experiments have been encouraging.

The gas electrode has allowed titration currents

of up to 50/u,A to be used without excessive polarisation, and subsequent equihbration times have

been reduced to 1-3 hours (cf. solid Ti02-x)

electrodes. However, at compositions approaching

stoichiometry {x in Ti02 - j < 0.001), and at tempera290-621 0-68— 12 169
tures in the range 900 to 1100 °C, the recorded

e.m.f.'s exhibited a continual drift with time

(1-5 mV per hour) which indicated that the experimental assembly was not sufficiently impermeable

to prevent changes in the oxygen content of the

gaseous environment of the non-stoichiometric

powder.

Two other cell assemblies are depicted in figure

6. Type V has been used by Tretyakov [27] and

Flengas [25] to measure equilibrium oxygen chemical potentials of various metal/ metal oxide systems.

Using a similar configuration it should be possible

to vary the composition of a powdered non-stoichiometric phase contained within the solid oxide

electrolyte tube. In practice, however, oxygen

permeabihty through the electrolyte tube could

be a major problem.

Rapp [29] has designed the type VI cell assembly

to determine the solubihty and diffusion coeffi- cients of oxygen in solid copper, and the same

arrangement can in principle be used to study the

thermodynamic behaviour of any solid variable

composition phase. The configuration has the ad- vantage of allowing oxygen to be titrated across both faces of the pellet and at the same time minimising the surface area of the pellet in contact with the surrounding gaseous phase.

3. Results

, :.

... 3.1.

Using a coulometric technique Markin [8] deter- mined the partial molar free energies of solution

of oxygen (ACqJ in UO2 + J as a function of composition. The reported values are in general agreement

with gas equilibration studies and have been the

basis for many theoretical models of the defect

structure of UO2+J [30] . The high mobility of oxygen

(Do,- = 10-»-10i« cm^ sec-i at 1000 °C [31], within

1!Electrometer

1 3 2 3 1

1 Non- stoichiometric electrode 2 Pt- O2 electrode

1 Working and Reference electrodes

2 Variable composition electrode

3 Electrolyte material

Figure 6. Cell assembly for solid state coulometric titrations {Types V and VI).

the nonstoichiometric UO2+X phase undoubtedly

contributed to this successful application of the

coulometric technique. It would be interesting,

however, to measure the partial molar heat of solu- tion of oxygen (AHq^) at compositions much closer

to stoichiometry than those investigated by Markin

and his co-workers [8] in an attempt to establish

the range of composition for which the simple

Wagner-Schottky model is applicable (constant

A//02). These studies would be particularly valuable

in view of the A//02 values recently determined

calorimetricaUy by Dode et al. [32]. These coulorimetric values are not in agreement with those

derived from the coulometric studies, and additional experiments are required to clarify this uncertainty.

3.2. Ti02-x

Experimental problems associated with cou- lometric studies on this oxide have already been

mentioned in preceding sections, and the results have also been examined in some detail elsewhere

[9]. These results are compared in figure 7 with

other data obtained using thermogravimetric tech- niques [33, 34] and the logarithmic plot indicates some difference in the values of the pressure de- pendence of the defect concentration. The slopes

of the isotherms, however, are very sensitive to small changes in the reported composition values

which are subject to errors of at least 0.001 in the

values of x.

It is more interesting to compare the A//02

values reported as a function of composition in figure 8. The agreement (A//o2=108±5 kcal)

between values obtained in using a thermogravimetric [34] , conductivity [35] and coulometric

[9] technique is encouraging. Also taking into consideration the different origins of the various

Ti02-j- samples it does appear that the A//02

values refer to the intrinsic defect properties of

rutile and are not determined by the presence of

aliovalent cation impurities. Also noteworthy is the

comparatively large composition range extending

up to a defect concentration of approximately

3 X 10-3 at 1000 °C for which the simple SchottkyWagner model appears to be applicable (constant

A//02). At higher concentrations the partial molar

heat term is a function of composition indicating

that interaction between the defects is becoming

important.

3.3. Nb02-x

Attempts to control the composition of Nb02-x

using the coulometric t-^chnique and the type II cell assembly (fig. 4) were unsuccessful [9]. Although it

170
Figure 7. Comparison of results at 1000 °Cfor the variation ofoxygen partial pressure with composition ofTiOs -x using a

logarithmic plot.

-60

AH

2^2

(kcol )

-70

-80

-90

-100

-110

-120

-130 -

-140

"I r

Nb02_x ( 9 )

TiOp.2-x

-H(34)

z3z~-_-_-_-.-_-^ (35 ) ,^19) ^

Hitit>tt^tx—tx—^ e ^

<(33 )

10 10- COMPOSITION X IN MOj.^

Figure 8. Comparison of partial molar heats of solution of oxygen as a function of composition for TiOi-^ and NbOz.^.

appeared feasible to add oxygen quantitatively to the oxygen-deficient Nb02, it was found impossible

to reverse the process and remove oxygen quantitatively. The slow diffusion process within Nb02-x

and the low oxygen chemical potentials involved

(- 100 kcal > A)Li02 (1000 °C) > - 130 kcal) probably

both contributed to the failure of the technique on

this occasion. The dependence of composition on

the oxygen chemical potential was therefore deter- mined [9] by equihbrating Nb02 pellets in fixed

PH2/PH2O gas mixtures, followed by emf measure- ments and gravimetric analyses. These data pro- duced a constant value for A//1/202 (~ 65 kcals) over

a wide range of composition (fig. 8), and long annealing times (~ 100 hrs) for compositions exhibiting

large departures from stoichiometry (x ~ 0.01) failed

to produce any evidence for ordering processes

due to the absence of any super lattice lines ^-ray

examination), or changes in the values for AHu-zo^-

It would be interesting to examine whether the

composition of this phase could be controlled by

coulometric titrations using the modified Type IV

cell assembly (fig. 5).

3.4. VO2-X and M0O2-X

Problems associated with the fabrication of

dence strong pellets of M0O2-J and V02-j- have

necessitated that coulometric studies [36] on the

thermodynamic properties of these phases be

conducted with powdered samples suspended from

an electrical microbalance in the type III cell, configuration (fig. 5). As reported earlier (section

2.3), however, it proved impossible to adequately

control the oxygen content of the gas phase due to the permeability of the zirconia electrolyte tube

and these studies have so far been inconclusive.

4. Conclusions

The present survey indicates that solid state coulometric titrations can provide a convenient

171
and reliable method for investigating the thermodynamics of nonstoichiometric phases possessing

high ionic mobilities and equilibrium oxygen

partial pressures in the range, 10"** > Pq atm

(1000 °C)>10-'6 atm (e.g., + X and Feb, + ^). However, slow solid state diffusion processes combined with relatively fast reactions between nonstoichiometric electrode materials and the sur- rounding gaseous environment at present severely

limit the application of this teclinique. Further

developments of the method are probably dependent upon the design of alternative cell assemblies

which are able to maintain the appropriate equilibrium oxygen partial pressure above the nonstoichiometric phase, and thus allow the relevant

electrode system at any given temperature and

composition to behave as an invariant system.

The comments on the modified cell designs are based on experimental work performed by Miss

S. Zador and Mr. J. Chan of the Metallurgy Department, Imperial College.

5. References

[1] C. Wagner, Thermodynamics of Alloys (Publ. by Addison- Wesley. Reading, Mass. (1952).

[2] J. S. Anderson and B. G. Hyde, J. Phys. Chem. Solids, 1967, 28, 1393-1408.

[3] S. Kachi and P. Roy, Phase Equilibrium Studies and

Transitions in the System V2O3-V2O5. presented at the 67th Annual Meeting, American Ceramic Society, 15th May, 1965 (Basic Science Division No. 80-B-65).

[4] A. S. Khan, Intermediate Oxides of Vanadium and Ti- tanium, D. Phil. Thesis, Oxford University, November

1967.

[5] C. Wagner, J. Chem. Phys., 21, 1819 (1953).

[6] J. B. Wagner and C. Wagner, J. Chem. Phys., 1957, 26,

1602.

[7] K. Kiukkola and C. Wagner, J. Electrochem. Soc, 1957, 104 , 379-386.

[8] T. C. Markin and R. J. Bones, United Kingdom Atomic

Research Authority, Report No. 4178.

[9] C. B. Alcock, S. Zador and B. C. H. Steele. Proc. Brit. Ceram. Soc. 1967, No. 8, 231-245.

[10] H. F. Rizzo, R. S. Gordon, and I. B. Cutler, Determination

of Thermodynamic Properties in Single Phase Wiistite by Coulometric Titration in a High Temperature Galvanic

Cell, presented at the National Bureau of Standards Conference on 'Mass Transport in Oxides' (24th October

1967).

[11] B. C. H. Steele and C. B. Alcock. Trans. A.I.M.E. 1965, 233, 1359-67.

[12] C. B. Alcock and S. Zador, Electrochim. Acta, 1967, 12,

673-7.

[13] B. C. H. Steele in Electromotive Force Measurements in High Temperature Systems, to be published by the Inst, of Mining and Metallurgy. London, April, 1968.

[14] P. W. Patterson, E. C. Bogren, and R. A. Rapp. J. Electro- chem. Soc. 1967, 114, 7.52-758.

[15] E. W. Roberts, Permeability Studies on Refractory Oxide

Systems, Ph.D. thesis, Leeds University U.K. (Sep- tember, 1966).

[16] R. Fabre, Etude de la Permeabilite aux Gaz des Ceramiques

d'oxydes a Haute Temperature, doctoral thesis, Grenoble

University, France (July, 1967).

[17] H. Smith, Metallurgy Department, Imperial College,

private communication.

[18] S. V. Karpachev, A. T. Filyayev, and S. F. Palguev, Elec- trochim, Acta., 9, 1681-5 (1964).

[19] C. B. Alcock and T. N. Belford, Trans. Faraday Soc. 1964, 60,822-835.

[20] C. B. Alcock, and T. N. Belford, Trans. Faraday Soc.

1965. 61, 443-453.

[21] G. P. Stavropoulos, Physicochemical Studies Associated With the Electrochemical Measurement of the Solubility of Oxygen in Liquid Sodium, Ph.D. thesis, London University, June 1967.

[22] J. S. Choi and W. J. Moore, J. Phys. Chem. 1962, 66,

1308-11.

[23] B. C. H. Steele, Thermodynamic Studies Using Galvanic

Cells With Solid Electrolytes, Ph.D. thesis, London University, July 1965.

[24] L. Himmel, R. F. Mehl and C. E. Birchenall, Trans.

A.I.M.E., 1953, 197, 827-843.

[25] R. Haul and G. Dumbgen, J. Phys. Chem. Solids, 1965, 26, 1-10.

[26] A. Kubik and C. B. Alcock, Metal Sci. J., 1967, 1, 19-24.

[27] Y. D. Tretyakov, Inorg. Materials, 1966, 2, 432-6.

[28] G. G. Charette and S. N. Flengas, Dept. of Metallurgy and Materials Science, University of Toronto, Canada,

submitted for publication in J. Electrochem. Soc.

[29] R. A. Rapp, Mixed Conduction in Solid Oxide Electrolytes, presented at the I.A.E.A. Conference in Vienna.

Sept. 1967.

[30] L. E. J. Roberts and T. L. Markin, Proc. Brit. Ceram. Soc. No. 8, 201-217 (1967).

[31] A. B. Auskern and J. Belle, J. Chem. Phys. 1958, 28, 171.

[32] M. Dode and P. Gerdanian, Comptes Rendues 1966, 2 62,

796-799.

[33] P. Kofstad, J. Phys. Chem. SoUds, 1962, 23, 1579-1586.

[34] K. Forland, Acta Chem. Scand., 1964, 18, 1267.

[35] R. N. Blumenthal, J. Coburn, J. Baukus, and W. M. Hirthe,

J. Phys. Chem. Solids 1966, 27, 643-654.

[36] S. Zador, Metallurgy Department, Imperial College, Private Communication.

[37] L. A. Simpson and R. E. Carter, J. Am. Ceram. Soc. 1966. 49, 139-144.

[38] W. H. Rhodes and R. E. Carter, J. Am. Ceram. Soc. 1966, 49, 244-249.

[39] H. S. Edwards, A. F. Rosenburg and J. T. Bittel, Tech.

Doc. Report No. A5D-63-635 (1963), General Electric Co.

[40] R. J. Hawkins, Cation Diffusion in Fluorite Lattices, Ph. D.

thesis, London University, Feb. 1967.

172
Interdiffusion Coefficients From Electronic Conductivity Measurements

Application to CugO

R. H. Campbell, W. J. Kass and M. O'Keeflfe

Chemistry Department, Arizona State University, Tempe, Arizona 85281

The interdiffusion of oxygen and cuprous oxide has been followed by monitoring the electrical conductivity of single crystal cuprous oxide specimens. The measured values are in good agreement

with the values calculated from the tracer diffusion coefficient of copper and the departure from

stoichiometry using the Darken equation.

Key Words: Cuprous oxide, electrical conductivity, interdiffusion, nonstoichiometry.

Introduction

In many non-stoichiometric oxides with ionized

vacancies, the conductivity is proportional to the

departure from the ideal composition. This is the case for example in CoO [1].' In CU2O the conductivity is accurately proportional to the square

root of the copper vacancy concentration (as shown

by the oxygen pressure dependence of conductivity

[2] and departure from stoichiometry [3]). It would

appear possible therefore to use conductivity as a measure of composition of the sample, and indeed

several authors [4] have used this fact to foUow the

kinetics of interdiffusion in oxides.

It is necessary however to proceed with caution

as in general the apparent conductance of an in- homogeneous sample is not proportional to the

average concentration of defects. Provided that there is local electrical neutrality (i.e., in the

absence of a space charge), it is possible however,

to relate the local conductivity cr to the local concentration of defects, c.

Theory

We have considered the case of a rectangular

prism — a\ < x\< a\,— a^< x%< a-i, — az< X'i<

with porous electrodes on the x\. xi faces so that the conductance can be measured in the direction. We suppose that the initial conductivity cr, is due to an initial homogeneous concentration ot defects c,. For times greater than zero the surface

concentration is held at c/ so that the conductivity comes to a value cr/ as t—^ =0. The concentration

at any point in the material is then given by the

well-known expression

X2, JC3) = C/-(C/-Ci)S(^,)S(^2)S(^3), (1)

where

5(^,) = (2/7r) I; [(-l)'V(n + l/2)] n=0

exp [-)8i(n + 1/2)^772] cos [(;; + l/2)7r^,] , (2)

and in turn

^i^xi/ar, fii = Dt/af. (3)

The local conductivity is cr(xi, x-z, X3) and as discussed above we have

Cr(xi, X2, X3) = A[c{Xu X2. Xs)]'", (4)

with A a constant and m usually 1 or V2. The apparent conductivity, cr, measured in the

X3 direction is obtained by multiplying the con- ductance of the sample by the length/area ratio and is related to the local conductivity by

03

2 ai 02

' r (13

I r

J -'I ,

dX3

C

(11 r (12

I I crdx-ydx^

J - (I, J -n.

(5)

From eqs (5), (4) and (1) one obtains after some

manipulation

1

' Figures in brackets indicate the literature references al the end of this paper.

(Tf—cr

(Tf—(Ti I — a -(/: (6)

173
where f f

' [1 l)5(fi)S(6)S(^3)]'"

and a = (7ilaf.

It is of interest to compare 2 as given by eq (6) with

the corresponding quantity W = (wj— w)l{wf — Wi)

where Wf and Wi are the final and initial amounts

of defects. The explicit expression for W is given

in ref. [4]. In figure 1 we show W and 2 as a function of the

dimensionless quantity Dtja'l for various shape

samples and values of ay/cri for the case m = \. It may be seen that in general W axvA X are very close

to each other for long, thin samples when oy/cr, is not very different from unity. In other cases however the assumption that 1,= W may lead to large errors in the calculated diffusion coefficients.

Application to Cuprous Oxide

Attempts to study interdiffusion in cuprous oxide

using a vacuum microbalance technique were

severely limited by the nature of the material.

SpecificaDy, at temperatures above 1000 °C evaporation is very rapid so that the small weight changes

upon change of stoichiometry are hard to deter- mine. At lower temperatures the cuprite phase

field is very narrow, again precluding reliable measurements. For these reasons we decided to resort to electrical measurements in this system.

Two single crystal specimens were used; sample

1 was 4.50 X 4.65 X 6.34 mm and sample 2 was

1.91 X 1.91 X 15.9 mm. In each case the conductivity was measured along the longest direction using

two probes. After equilibration at low oxygen pressure, oxygen was rapidly admitted to the sample

and the conductivity monitored as a function of

time. Strictly, in order to use the expressions

derived above for 2, porous contacts covering the

whole face of the crystal should be used. The use

of smaller contacts is hkely to be less serious in the case of sample 2. As conductivity changes are very rapid, the sample current with known constant

applied voltage was recorded with a recorder speed

of 6 inches/minute. Plots of 2 versus log t were

superimposed on theoretical plots of 2 versus log

{Dtla'D calculated according to eq (6). The dif- ference in the abscissa of the two plots is equal to log {Dial) so that D is readily calculated. As the

conductivity is very temperature sensitive it is quite possible to measure erroneous values of oy; however, this becomes evident when the experimental curve is compared to the theoretical curve, and any experimental results differing significantly

from the theory were rejected.

-3 -z -1 o 1

-+ -3 -2 -1 o

1.0

O.S

1 1 —

w ^

1 1 1 -6 -S -4 -3 -Z log (Dt/a|)

Figure 1. Thereotical curves for conductivity and weight change

as a function oftime for various shape samples.

The numbers on the curves are the values of o-j/oy. The curves labeled W refers

to weight change, the others to conductivity change.

Because of the very rapid diffusion observed

there was considerable scatter of the experimental

results. These are reported in table 1 with results

for each crystal at each temperature reported as an average. The results for sample 2 were in general better and are in fair agreement with the

theoretical values (discussed below). The results

for sample 1 are too high; it is not clear at present

whether this is due to the fact that the experimental conditions did not conform to the boundary

conditions used in deriving eq. (6) or whether ef- fects such as surface diffusion are significant.

Theory for Cuprous Oxide and Comparison

With Experiment

We can calculate the expected interdiffusion

coefficient as follows. Neglecting oxygen diffusion,

which has been shown to be very much slower than

copper diffusion [5], the rate of interdiffusion of

174
oxygen and CU2O1 + J is given by the Darken

equation [6]

5=(l//)D*„(dlnacu/dlnccu), (7)

where D is the measured interdifFusion coefficient,

D*^ is the tracer diffusion coefficient and / is a correlation factor to take into account the fact that successive jumps of the tracer atom will not be random. For a face-centered cubic lattice (as in the case for Cu+ in Cu2O)/=0.78. In the final factor in eq. (7) Ocu is the activity of copper in cuprous oxide and Ccu the concentration of copper.

According to Moore and Selikson [7] at an oxygen pressure of 0.1 torr

log,oZ>*u = -l-36-7900/r. (8) We have also that

where Oo is the activity of oxygen and to an excellent approximation acu20 is constant, so that

dlnacu = -(V2)dlnao (9)

or, as Oo is proportional to the square root of the

oxygen pressure,

dhiacu = -(V4)dln Po^ (10)

In terms of the stoichiometric excess of oxygen,

X, we have also for the concentration of copper,

measured in fraction of sites occupied,

Ccu-i2-x)l2 (11)

or, recalling that x <^ I,

d\nccu = -dxl2 (12)

so that finally

dlnacu/dlnccu = (V2x)(dln PJdlnx) (13)

O'Keeffe and Moore [3] found by direct gravimetry

that at f 02 = 0.1 torr

login ^ = 0.02 -4760/r (14)

and (din PJdlnx) = 3.7 (15)

Substituting into eq. (7) the appropriate values

from eqs. (8), (13), (14) and (15), we obtain finally

for the calculated rate of interdiffusion

log,oD= 1.00- 3140/r (16)

It should be remarked at this point that the only

assumptions made about the mechanism of diffu- sion are (a) that oxygen diffusion is negligibly slow

and (b) that the correlation factor is 0.78. In particular it should be emphasized that if there is a

diffusion potential and there is ambipolar diffusion

of mobile holes and less mobile vacancies this factor

is already included in the term (din acu/dln Cqu) in eq. (7). This point has been discussed at some length

elsewhere [8]; the general validity of eq. (13) is illustrated by a consideration of two specific models

in the appendix.

We turn now to a comparison of the experimental

results with the theoretical expression, eq (16). In

table 1 we have recorded the experimental interdif- fusion coefficients for both samples at a number of

temperatures. Also in the table are the diffusion coefficients calculated from eq (16). For sample 2 the

agreement with the calculated values is very good;

sample 1 gives values that are consistently too high.

We beUeve that the discrepancy for sample 1 arises

because the theoretical conductivity curves were

calculated for contacts covering the whole face of the crystal, whereas in fact "point" contacts were

used. As the area of the contact faces was small

in comparison to the area of the other faces for sample 2, the error should not be so great for this sample.

The agreement between theory and experiment

for sample 2 is very gratifying when it is recalled

that the results of three separate measurements

(interdiffusion, tracer diffusion and nonstoichiometry) are being compared. The most difficult meas- urement is that of the degree of nonstoichiometry,

x, and we consider the present measurements to be a valuable confirmation of the order of magnitude at least of the previously reported results. The D reported here may be interpreted as a vacancy diffusion coefficient. The very high difTable 1. Measured and calculated values of logjo D in CusO.

T °C Sample 1 ^ Sample 2

" Calculated

933 ±1 -3.4 (1) -3.6 (2) -3.60

973 ±1 -3.15 (5) -3.4 (2) -3.52

1001 ± 1 -3.2 (3) -3.75 (4) -3.46

1032+1 -2.6 (2) -3.45 (3) -3.40

1064+1 -2.5 (2) -3.4 (5) -3.36

^Numbers in parentheses are the number of separate measurements at that temperature. Calculated from eq (16).

175
fusion coefficients are in accord with the observation that internal equilibrium can readily be reached

in cuprous oxide down to very low temperatures

[2, 9|. It is worth remarking also that in general

cation mobility is very high in the chalcogenides of

the univalent metals.

This research was supported by the Air Force

Office of Scientific Research, Office of Aerospace

Research, United States Air Force, under AFOSR

grant Nr. AF-AFOSR-717-65.

Appendix

We consider here more fully the question of a

diffusion potential and demonstrate that the vaHdity

of eq (13) is not dependent on a particular model

for the incorporation of oxygen in cuprous oxide.

Two possible models are considered:

(a) production of neutral cation vacancies

crystal^ Fcu+ Cu(Has); (Ala)

(b) production of charged (ionized) vacancies

crystal-^ F2- + ne + +Cu(j,as). (Alb)

In either case, assuming as before that the activity

of the crystal is constant, and assuming local thermodynamic equilibrium, we have the equilibrium constants:

acua(Vcu)=Ka. (A2a)

acna{V'^-)aHe^) = K,. (A2b)

Here the activities, «, are defined in terms of the

electrochemical potential fx, by

a-exp[(M-At°)//?rj. (A3)

In a dilute solution of ions of charge Z

a = c exp{Ze(t)lkT), (A4)

where Z is the valence of the ion, c is the concentration, measured most conveniently as site fraction,

and 4> is the local electrostatic potential. Thus

a(Vcu) = xl2,

a{V^-) = {xl2) exp i-necjilkT),

a(e + )-(nx/2) exp (e(^)//.T).

The equilibrium constants of eqs (A2) now become

acu{xl2)^ Ka, (A5a)

acJxl2)inxl2}" = Kb. (A5b)

It should be noted that the local electrostatic potential does not appear in these expressions.

It is now possible to evaluate the terms in eq (13); as previously we restrict ourselves to the situation

where <^ 1 so that In Cqu ~ (— jc/2), and obtain:

Case (a) din acu/dln Ccu= (2/:k); (A6a)

Case (b) din Ocu/dln Ccu= [2(n+ 1)/^]. (A6b)

On the other hand using eq (10) it may be seen

that from eqs (A5):

Case (a) din Po^/dln x = i; (A7a)

Case (b) din PJd\n =i4 ( /i + 1 ) . ( A7b)

In both cases eqs (A6) and (A7) yield

din ocu/dln Ccu= (l/2;c)(dln PoJd\n x), (13)

as we set out to demonstrate.

Finally it may be remarked that the result, eq (15), quoted indicates strongly that case (a) obtains in cuprous oxide (cf. eq (A7a)). Recent high temperature Hall-effect measurements [10] confirm this observation.

References

[1] B. Fisher and D. S. Tannhauser, J. Electrochem. Soc, 111,

1194 (1964). [2] R. S. Toth, R. Kilkson and D. Trivich, Phys. Rev., 122,

482 (1961); M. O'Keeffe and W. J. Moore, J. Chem.

Phys., 35, 1324 (1961). [3] M. O'Keeffe and W. J. Moore, J. Chem. Phys.. 36, 3009

(1962). [4] J. B. Price and J. B. Wagner, Z. phys. Chem. (N.F.), 49, 257

(1966) ; L. C. Walters and R. E. Grace, J. Phys. Chem.

Solids, 28, 245 (1967): D. G. Thomas, J. Phys. Chem.

Solids, 3, 229 (1957). [5] W. J. Moore, Y. Ebisuzaki and J. Sluss, J. Phys. Chem., 62,

1438 (1958); M. O'Keeffe, Y. Ebisuzaki and W. J. Moore,

J. Phys. Soc. Japan, 18, 131 (1963). [6] L. S. Darken, Trans. Met. Soc. AIME, 175, 184 (1948). [7] W. J. Moore and B. Selikson, J. Chem. Phys., 19, 1539

(1951); 20, 927 (1952). [8] M. O'Keeffe, Diffusion in Oxides and Sulfides in Sintering and Related Phenomena (G. C. Kuczynski, Ed.), Gordon

and Breach, N.Y. (in press). [9] M. O'Keeffe and F. S. Stone, Proc. Roy. Soc. (Lond.), A267 , 501 (1962).

[10] H. L. McKinzie and M. O'Keeffe, Phys. LeU. 24A, 137

(1967).

176
A Chemla Experiment in BeO ^

C. F. Cline, H. W. Newkirk, R. H. Gondii, and Y. Hashimoto

Lawrence Radiation Laboratory, University of California, Livermore, California 94550

The drift of Be' tracer in BeO at 1550 °C under the influence of an externally apphed d-c electric

field was determined by a sectioning technique. Reasonable agreement was obtained with the NernstEinstein relation.

Key Words: Diffusions, electriceJ conductivity, tracer diffusion.

The drift of a radioactive tracer under the in- fluence of a d-c electric field has been described

by Chemla [1]^ and referenced by others [2, 3], In this case, the diffusion equation for a cation

tracer is

dc _ d^c dc

'dt~^ dx^ dx' (1)

where C is the concentration of the tracer, D is the

diffusion coefficient, /x is the ion mobility of the

tracer, and E is the electric field strength. For an

anion, the term in E would be positive. The net

effect of the electric field is simply to displace the center of gravity of the diffusion distribution profile by a distance,

^= IxEt. (2)

There is no other distortion of the distribution profile if the field is homogeneous in the region of

diffusion. Experimentally, two flat slabs are pressed

together with a thin layer of radiotracer between

them, and a high-temperature anneal is carried out with the electric field imposed normal to the plane.

The resultant distribution of tracer can be analyzed

in terms of the distance, x, from the initial interface,

Mr 1 {x-^xV

where Co is the initial amount per unit area and

C(x) is the concentration at a distance, x. Thus, the

mobility can be found from an experimental deter- mination of the center of gravity of the profile since

the maximum of the beU-shaped profile moves

' Work performed under the auspices of the U.S. Atomic Energy Commission.

^ Figures in bracicets indicate the literature references at the end of this paper.

away from the interface with a velocity, fxE. The

diffusion coefficient can be obtained in the usual

way from the shape of the profile, i.e., from plot

of log C versus (x — Ax)^.

The work reported here was concerned with

measurement of the mobility of Be'^ tracer in BeO.

The material used was the same as that previously

employed in studies of self-diffusion in BeO [4].

It was 98 percent dense, contained grains averaging

30 microns, major impurities were Al (200 ppm) and

Si (50 ppm). BeO has the hexagonal, wurtzite crystal

structure, and there is a preferred crystaUographic

orientation among the grains; the majority of the c-axes of the crystalhtes were at right angles to the field direction. Three pieces were prepared

with front and back sides made parallel, the tracer was deposited on the polished faces by a technique

previously described [4] , and the individual pieces were assembled as shown in figure 1. The attach- ment of electrodes to the stack and the placement

of electrodes used to measure the potential drop

across each interface are also shown. Platinum

electrodes were used on the end samples, and a potential of 30 volts was supplied throughout the

run. The current passing through the assembly was monitored by a recording ammeter. The duration of the anneal was 5.15 X 10^ seconds and the

temperature 1550 °C. While there was a considerable voltage drop at the electrode BeO interface,

the field strength was only about 10.0 volts per

centimeter in the neighborhood of the tracer. Following the anneal, the distribution of tracer in the region between the samples 1-2 and 2-3 was

determined by sectioning. The resulting concentration profile is shown in figure 2. If the tracer had behaved ideally, the profiles in figure 2 would follow eq (3) and have a relatively

smooth, displaced bell shape. However, there is some barrier to completely free interchange of

177
POWER

VOLT

METER

-<!>-

POWER SUPPLY - REGULATED HIGH-VOLTAGE SUPPLY MODEL 241 KEITHLEY INSTRUMENTS VOLT METER — DIFFERENTIAL VOLTMETER MODEL 801 JOHN FLUKE, SEATTLE AMMETER - ESTERLINE-ANGUS GRAPHIC AMMETER MODEL AW

0-1.0 dc mA

Figure 1. Circuit diagram.

tracer between pieces of BeO. One also sees an

apparent decrease in activity concentration near

the surface, but this is probably due to slight pitting

of the specimens when they are pulled apart after the anneal. The tails of the profiles are enhanced

by grain boundary diffusion. Since the 1-2 interface exhibited much greater

surface holdup (fig. 2) than the 2-3 interface, the

latter was used for all calculations. The displacement of the center of gravity of the diffusion profile can be estimated as 20 microns for interface 2-3.

In both cases, the drift of the tracer was toward the

cathode. Using this displacement and a field gra- dient of 10 V/cm, the mobility calculated from eq

(2) is found to be 3.9X10-''' cm'^-volt-'-sec"'

(+) ANODE END

100 50 0 50 100 150

CALCULATED MEASURED

0 50 INTERFACE DISTANCE — microns

Figure 2. Be' tracer distribution profiles.

It now becomes possible to calculate the dif- fusion coefficient for Be+^ ions from the mobility

through use of the Nernst-Einstein relationship.

fxRT

zF (4)

where R is Boltzmann's constant, 8.31 joules deg~^

mole-\ F is a Faraday, 9.65 X lO* coulomb mole"'

and Z, the charge per ion, is 2. One finds at the

2-3 interface, A = 3.05xl0-" cm^-sec"!.

The diffusion coefficient as obtained by tracer techniques, Dsd, is related to that obtained from

ionic transport measurements, Du by the correlation coefficient [2, 3],

f-D,,IDi (5)

Using 0.78 for / [5, 6, 7], one calculates a Dgd of

2.38x10"" cm^-sec"^ Substituting the time and

Dsd back into eq (3), one can calculate the diffusion

profile. This calculated curve is shown in figure 2, and it appears to agree fairly well with the measured points. The diffusion coefficient calculated

here is only about 40 percent of the value to be

expected on the basis of previous measurements

[4]. But this is not a serious disagreement in view

of the scatter in the data. In conclusion, it has been found that Be tracer migrates toward the cathode at a rate which tends

to support the view that electrical conductivity in BeO is by a mechanism involving beryllium ion

diffusion.

The authors are indebted to Prof. C. E. BirchenaU and Drs. L. Himmel and R. Borg for reviewing

the manuscript.

References

[1] M. Chemla, Diffusion of Radioactive Ions in Crystals, Ann.

Phys. (Paris) ser. 13, 1, 959. Also UCRL-Trans 1038(L)

(1956).

[2] A. B. Lidiard, Ionic Conductivity, Handbuch der Physik

20, 246 (1957).

[3. J. R. Manning, Isotope Effect in an Electric Field and Jump

Frequencies for Diffusion in Ionic Crystals, J. Appl. Phys.

33, 2145 (1962).

[4J R. H. Condit and Y. Hashimoto, Self-Diffusion of BeryUium

in Beryllium Oxide, J. Am. Ceram. Soc. 50, 425 (1967).

[5] K. Compaan and Y. Haven, Correlation Factors for Diffusion

in Solids, Part 2 — Indirect Interstitial Mechanism, Trans. Faraday Soc. 54, 1498 (1958).

[6] H. J. de Bruin, G. M. Watson, and C. M. Blood, Cation Self- Diffusion and Electrical Conductivity in Polycrystalline

BeO, J. Appl. Phys. 37, 4543 (1966).

[7] C. F. Cline and H. W. Newkirk, Electrical Transport Proc- esses in Beryllium Oxide, to be published.

178
Self-Diffusion of Oxygen in Neodymium and Samarium Sesquioxide

G. D. Stone, G. R. Weber and L. Eyring

Department of Chemistry, Arizona State University, Tempe, Arizona 85281

The oxygen self-diffusion coefficients were measured for Nd203 (^-type) and SmzOs (B-type) and

an estimate was made for the diffusion coefficient of Er203 (C-type) at 1000 °C using the method of isotope exchange. The rate of change of '^0 in the gas phase in equiUbrium with a solid sample was

measured by a mass spectrometer. The enriched oxygen gas was circulated over the sample and through

a leak valve which continuously admitted a small amount of the gas directly into the analyzer of the mass spectrometer. The solid samples were in the form of spheres prepared in an argon plasma. Over

the range of pressures used (30 to 300 mm Hg), there was no pressure dependence noticed. The self- diffusion coefficients as a function of temperature, determined in the range of 700 to 1000 °C using an

exact solution of Kick's second law with a first order exchange reaction at the surface, were

SmzOa (99.9%), Z) = 9.2 X 10-« exp(-23.5 ±2 kcalIRT) SmzOs (99.998%), Z) = 6.0 x IQ-" exp(-21.3 ± .5 kcal/«7') NdzOa (99.999%), Z) = 1.3 x lO"" exp(-31. ±2. kcal/RD

The value for the diffusion coefficient of Er203, estimated at 1000 °C, was 10"''' cm^/sec. The low

value measured for the heat of activation and the lack of pressure dependence indicated that the

diffusion was probably extrinsic. While the systematic errors were large, the measured values for the heats of activation indicated a possible correlation to structure and the lanthanide contraction.

Key Words: Oxygen isotope exchange, rare-earth oxides, self diffusion.

1. Introduction

In the study of the rare earth oxides, there have

been many indications that oxygen is much more

mobile than the cations. For example, in praseo- dymium and some of the other rare earth oxides

of variable valence, tensimetric studies have shown

that the soUd reaches equiUbrium with oxygen

very rapidly after a change in the oxygen gas pressure. This has been observed as low as 400 °C; far below the Tammann temperature where the cations

should be mobile [1, 2].' In the oxidation of praseo- dymium metal, scratches on the surface appeared

unchanged after the formation of a dense oxide

layer upon complete oxidation, indicating that oxygen was the mobile species [3]. Oxygen mobihty

has also been shown to be important in the vacuum

reduction of PrOi 71 where the self-diffusion coefficients for oxygen for particles 0.5 microns in diameter have been estimated to be [4] D=1.5XlO-« cm2/sec at 990 °C

D = 1.7 X 10-" cm^/sec at 700 °C.

' Figures in brackets indicate the literature references at the end of this paper.

In a more recent study on the oxidation of rare earth metals, an estimate of lO"*^ cm'^/sec at 1600 °C was made for the self-diffusion coefficient of oxygen

through the oxide coating [5]. Using shghtly re- duced Y2O3 (C-type) and EraOs (C-type), diffusion

coefficients for oxygen were recently given as a function of temperature by the relations [6]

Y2O3, D = 7.24 exp(-58.6 kcal/RT),

for 1000 to 1500 °C

EraOs, D = 1.22 exp(-47.8 kcal/RT),

for 850 to 1250 °C.

All of these measurements indicated that the

self-diffusion coefficient of oxygen in the rare earth

oxides appeared to be much greater than those

observed for some other transition metal oxides;

for example, TiOa has a value about lO-'^cm^/sec

at 1000 °C [7]. In order to improve understanding of the apparent

high mobility of oxygen, it was desirable to measure

the self-diffusion of oxygen and its temperature

and pressure dependence by a more direct method.

The method chosen for the study was that of

heterogeneous isotope exchange. This study was

179
of fixed valence oxides to see if a correlation could

be made between the diffusion coefficients and

structure. However, the apparatus was designed

to study either fixed or variable valence materials.

Since the high temperatures used in the preparation of the samples and the conditions of the

exchange experiments did not allow the use of one

cation for the three major polymorphic forms of

the sesquioxide {A, fi, and C-type), three different rare earths were used: Nd2 03 (^-type), Sm203

(fi-type), and EnO^ (C-type).

The structures of the A, B, and C-form rare earth

oxides may be viewed as being derived from the

fluorite structure characteristic of the dioxides of

these elements. In the fluorite structure each metal

atom is coordinated with eight oxygen atoms at the corners of a cube with all coordination cubes

sharing edges. Alternatively one may view the

structure as tetrahedrally coordinated oxygen

atoms in which all tetrahedra share edges.

In terms of MOn coordination cubes the C-type

structure is seen as six coordinated M0(;V2 (Vsymbolizes an oxygen vacancy) cubes with the two

vacancies occurring either on the face (in 75 per- cent of the cases) or body diagonals. The A and

Z?-type are seen as structures in which distorted

M0« coordination cubes provide the correct stoi- chiometry by shearing to allow face sharing as well

as edge sharing as has been illustrated by Eyring

and Holmberg [1]. Actually changes in M-0 dis- tances suggest seven coordination for A type and

six and seven for fi-type.

From the alternate viewpoint Caro [8J has illus- trated how these structures may be considered as three dimensional arrangements of OM^ tetrahedra

in which sheets of (MO);j+ polymers are seen to be

separated by oxygen layers. In either view the

oxygens are not all equivalent and one would expect

the oxygen transport properties to be determined

by the mechanism of transport which would be

sensitive to the structures and the interatomic

distances.

2, Apparatus

A block diagram of the system used is shown

in figure 1. The total system included gas preparaBa lance and _ SajTi^le

Heat Sink

Exchange

Somple

Pressure Measu rement

Cylinder Oxygen

Vacuum

Pump Oxygen-18

-®- o Isotope

Equilibration

Circulating Pump

-| Leak

Valve

External Sample

Mass

Spectrometer

Figure 1. Apparatus for oxygen isotope exchange.

tion and storage, vacuum pumping and pressure

measurement, and exchange and analysis sections.

A more detailed drawing of the exchange and

analysis section is shown in figure 2. The enriched

oxygen was prepared by electrolysis of water en- riched with ^**0 using NaOH as an electrolyte. Any

carbon dioxide absorbed in the water was removed

by repeatedly freezing the hquid and evacuating

Heat Sink With Samples

Figure 2. Samples and gas flow.

180
the gas over the ice. Any hydrogen in the oxygen was removed by oxidation catalyzed by Pd coated

asbestos heated to 400 °C. The cyhnder oxygen,

containing the normal isotopic concentration of

1^0, 0.2 percent, and the enriched oxygen both

had the water removed by dry ice traps before

storage in evacuated flasks. Mass spectrometric

analysis of the enriched oxygen showed only slight

increase in impurities over that normally measured

as background in the mass spectrometer.

The system was designed so that the gas in all sections could be evacuated or its pressure measured. For gas pressures above 1 mm Hg, a mercury

manometer was used. A spiral gauge, used as a

null meter, isolated the system _ from mercury.

The system had been designed to determine the

composition as well as measure the rate of ex- change for the rare earths of variable valence.

For this reason two samples are shown in Figures

1, 2, and 3. Only one sample, the exchange sample,

3i ' CIRCULATING PUMP

.TO LEAK VALVE

INCONEL-702

THERMOCOUPLE

WELL

FOR SUPPORT RODS

The gas used for the isotope exchange was mixed

and circulated using a magnetically operated, in- ternal circulating pump over the sample and

through the leak valve as is shown in figures 3

and 4. The Granville-Phillips leak valve allowed

To Mass

Spectrometer ^^^{

Pyrex^

< Monel Bellows

Leak Valve Figure 4. Gas sampling.

only a small amount of the gas to pass continuously

into the analyzer section of the CEC 21-130 mass

spectrometer. Even after twenty-four hours, the

amount of gas lost from the system through the

leak valve was less than could be measured as a pressure change. Separate experiments using

nitrogen in one section mixed at a certain instant

with oxygen in the other established the time re- quired for gas mixing by the pump and indicated

the lack of serious memory effects in the con- tinuous sampling technique.

It took 3.5 seconds for the gas to go from the

sample to the leak valve, hence all measured times

had to be corrected by this amount. Jiru and

Novakova [9, lOJ found that the equilibrium concentration of the isotopes in the various molecules

of oxygen, i60'«0, isQi^O, and '«0i«0, was not always maintained. Therefore, a heated Pt + 10

percent Rh wire was used as a catalyst to see if this would cause a problem for the rare earth

oxides. It was found that equilibrium was estabhshed in all runs reported here since the temperature of the sample container was above 700 °C. A

run made at 1000 °C without a sample indicated

that no correction was necessary under the conditions of the experiments.

Figure 3. Samples and heat sink.

was used for the materials of fixed valence: Sm203,

Nd203, and Er-^Os. The furnace was designed,

with top and bottom auxiliary heaters and an

Inconel metal block, so that there would be a six inch constant temperature zone. The temperature

gradient of this zone was measured by difference

thermocouples (Pt/Pt + 10 percent Rh) inside

the furnace next to the furnace tube. The measured gradient was always kept within 1°. The temperature itself was controlled electronically to within 0.1° at 1000 °C with the temperature being

measured at the Inconel block with a Pt/Pt+10

percent Rh thermocouple.

3. Operating Procedure

Before an exchange study was carried out the

sample was heated in vacuum to 1000 °C in order

to remove adsorbed gases. After the experimental

temperature had been set, one of two procedures was used. One procedure, which would have been

appropriate for samples of variable valence, used

cylinder oxygen in the sample volume and enriched

oxygen in the rest of the system. In this case the

circulating pump mixed the gases as well as cir- culated them. In another procedure, the sample

volume was initially evacuated and enriched oxygen was admitted from the rest of the exchange system

at the beginning of the run. Tests on a sample of

181
Nd203 indicated that there was only a slight dif- ference between the diffusion coefficients measured

by the two methods. In either case, the amount of gas added was determined from the measured pres- sures and temperatures and the previously cahbrated volumes.

The i«0"'0+/'^Oi**0+ ratios were measured

with the mass spectrometer to determine the rate of change of concentration in the gas above

the exchanging sample. After an experiment on a sample at one temperature was finished (this

normally took one hour) cylinder oxygen was added

at 200 mm Hg to the evacuated exchange system.

The temperature was taken above 800 °C, and the

'^O was removed from the solid by back exchange.

The procedure was repeated four times with the

fourth batch of cyhnder oxygen being left overnight. The isotopic concentration of the solid was

determined from this last back exchange and in every case was 0.2 percent '^O. With this pro- cedure it was possible to use the same sample for a variety of runs.

4. Samples and Sample Preparation

The enriched water, used in the preparation of

the '^O enriched oxygen gas, was obtained from

Yeda Research and Development Co. and had a

listed enrichment of 1.66 percent '^'O and 0.14 per- cent ''O. The sohd samples, SmjOa and EraOa were obtained from Research Chemicals. The purity for two of the samples of Sm203 and for Er2 03 was

99.9 percent. The higher purity sample of Sm203

was 99.998 percent. Trona suppHed the 99.999

percent purity Nd203. All of the purities hsted are with respect to rare earths. The procedure for the preparation of spheres

from these starting materials consisted of two main

steps; sintering in vacuum at 1700 °C and melting

in an argon plasma. Before sintering it was usually

necessary to heat the fresh sample in a Pt crucible

in air to 900 °C to remove any hydroxide or car- bonate that had formed on standing. The samples

were also kept in argon or vacuum at every stage

of their preparation to prevent the readsorption

of H2O or CO2. Since not all of the sintered powder

that was put through the plasma melted to form

spheres, the spheres were separated by rolhng

them away from the odd shaped powder on a glass

plate held at a slight incline. The spheres were

then separated according to size by sieves. Microscopic examination of the whole spheres

at 50 X and of cut spheres at 400 X indicated that

5 percent were not melted throughout. The rest of

the spheres appeared to be homogeneous and trans- parent to translucent. Under a polarizing microscope, there was no evidence that the spheres were

single crystals. There was no appreciable change

in the appearance of the spheres after repeated

reaction and back exchange.

5. Computation

Solutions of Pick's second law for the isotopic

exchange of a solid with a well-stirred gas of limited

volume have been given by several authors including J. Crank [11] and R. Haul and coworkers [12,

13]. Preliminary results with the rare earth oxides

indicated that the simpler solution given by Crank

could not be used. It was necessary to consider a phase boundary reaction as well as diffusion in the

first portion of the reaction. A solution for this case has been obtained by Haul et al. [13]. Their solution

is:

/sVd ^\

where the Z„ are the roots of the equation

and Z4 = Zi , Z5 = Z2

.

The solution was obtained by making the assumption that in the subsidiary equation, given below,

cosh qa can be replaced by sinh qa.

P =

HP2-P,) aq cosh qa sinh qa

3K (^q cosh qa—- sinh qa + ^ smh qa

where q^ = ZID.

182
From the definition of q it can be seen that this is certainly valid when D is small enough. However,

the values measured in the present work were fairly

large {D ~ 10~^°). Therefore, it became necessary

to know how closely the approximate and exact

solutions agreed. Thus the diffusion equation was

solved using the method of the Laplace transform

and the boundary conditions given by Haul et al. [13]. The result is:

6. Results

6.1. Sin203 Experiments

Using 99.9 percent and 99.998 percent SmaOs,

three different samples were used to obtain the

diffusion coefficients and the phase boundary re- action constants listed in table 1.

Pi 1 + A 54^

A

-a-S-Dtl9 6Q\HlQHl + llk)

where the a„ are the roots of

9a„Q-af,\

where

The above equation reduces to that of Crank in the limit as K approaches infinity.

The agreement between the two solutions is good until the reaction is very nearly complete.

For the approximate solution the error function of a real variable was calculated using a rational

approximation [14] and the error function of a complex variable was calculated using an infinite series approximation [14]. When using the exact

solution twenty terms in the series were summed.

Both equations were used in the analysis of the

data. The diffusion coefficients obtained by the two

equations were essentially the same except in a few instances when the diffusion coefficients were

large. All calculations were performed on a CDC3400 computer which carried ten digits. The time

required for the calculations by the two equations was comparable.

When the logarithm of the diffusion coefficients were plotted against l/T. as shown in figure 5. two

Figure 5. Sm-^O^ Temperature dependence o/D.

separate hnes were obtained. One line corre- sponds to the 99.9 percent sample and the other

to the 99.998 percent purity sample. It can be seen

that the slopes are the same within the limits of error and only the intercepts are different. When

expressed in equation form:

Sm.Oa (99.9%),

£) = 9.2 X 10-6 (-23.5 ±2 kcsA/RT)

Sm-zOi (99.998%),

Z) = 6.0 X 10-6 exp (-21.3 ±.5 kcaVRT).

5.1. List of Symbols

P = isotope content in the gas at t = t

P-i = isotope content in the gas at t — o Pi — isotope content in the solid at t — o

rej,= moles of exchanging atoms in the gas

rig = moles of exchanging atoms in the soUd

k = nglns

D = Diffusion coefficient

S = specific surface = 3/radius of the sphere = phase boundary reaction constant

z = parameter of the Laplace integral

J^" e-"p{t)dt = p{z)

The variation with pressure, as shown in figure 6, indicates that there is no pressure dependence

X

M. X

•

•

—1

X —1

X

-i 1 1 1- 1—

X

1.5 1.7 1.9 2.3 2.5 LOG (P) (mm Hg) Figure 6. Sm203, Pressure dependence o/D.

183
Table 1. Samarium sesquioxide

Temperature

sample (°C) Pressure mm of Ha

D

cm'Vsec K cm/sec

Sample I: 1.3449 gm, 99.9% purity, spheres

0.1 11 ±0.038 mm diameter

Sm203 (99.9%), 0.17 ±0.1

SmaOs (99.998%), - 0.04 ±0.04.

6.2. NdoOs Experiments

Using 99.999 percent Nd203 spheres, three differ- ent samples were used to obtain the diffusion coeffi- 696.20 63.45 2.947 X 10-" 6.454 X 10-« cients and ph ase boundary reaction constants listed

794.10 899.50

64.20 64.10

1.237 X 10-'° 3.530 X 10-1" 4.150 X 10-^ 1.803 X 10-« in table 2. For the temperature dependence, two

793.10 119.85 1.135 X lO-i" 4.702 X 10-^ 793.10 214.10 9.000 X 10-" 5.926 X 10-' Table 2. Neodymium sesquioxide.

Sample II: 1.3980 gm, 99.9% purity, spheres

0.111 ±0.038 mm diameter Temperature Pressure D K

Sample (°C) mm of Hg cm^/sec cm/sec

907.60 44.80 2.542 X 10-'» 1.806 X lO-** 950.20

906.80

44.80

73.75

3.278 X 10-'» 2.740 X lO-'"

3.579 X 10-« 1.643 X 10-« Sample I: 1.2529 gm. 99.999% purity, spheres 0.160± 0.017 mm

diameter

907.50 104.70 3.125 X lO '" 2.542 X 10-« 907.60 145.40 2.197 X 10-'° 2.983 X 10-« 907.10 196.45 2.756 X 10-'° 2.116 X 10-6 900.20 36.30 1.023 X 10-^ 4.814 X 10-6 900.00

899.90

902.90

898.10

87.55

99.25

139.50 178.70

1.047 X lO-s* 7.810 X 10->» 9.694 X 10-'° 1.242 X 10-"1.004 X 10-5 9.137 X 10-6 1.113 X 10-» 1.235 X 10-5

Sample III: 0.8973 gm, 99.998% purity, 0.160±0.017 mm diameter

spheres

702.50 38.70 9.309 X 10" 2.287 X 10-' 897.60 286.05 1.129 X 10-" 3.829 X 10-^

725.50 38.75 1.435X10-'° 2.000 X 10-' 749.50

774.10

38.30

37.55

1.556 X 10 '° 2.196 X 10-'° 3.433 X 10-' 5.087 X 10-' Sample II: 0.99279 gm, 99.999% purity, spheres 0.163 ±0.014 mm

diameter

798.50 38.85 2.937 X 10-'° 8.545 X 10-' 822.90 38.85 3.430 X 10-'° 1.208 X 10-6 848.70 38.80 4.608 X 10 '° 1.725 X 10-° 706.80 74.75 1.741 X 10-" 1.629 X 10-' 874.90 38.95 5.742 X lO-'o 2.106 X 10-6 780.10 76.85 6.700 X 10-" 6.300 X 10-' 901.30 38.95 7.695 X 10-'° 2.777 X 10-6 852.90 78.30 1.740 X 10-'° 1.978 X 10-6 900.30 38.75 6.397 X 10-'° 2.762 X 10-6 924.20 77.05 3.701 X 10-'° 3.789 X 10-6 924.50 39.15 8.416 X 10-1° 4.819 X 10-6 1002.80 75.90 7.300 X 10-'° 6.641 X 10-6 953.50 38.90 9.491 X 10-'° 4.195 X 10-6 973.10 38.80 1.001 X 10-^ 6.535 X 10-6 998.70

799.90

799.40

38.60

50.10

76.15

1.221 X 10-» 2.250 X 10-'° 1.894 X 10-'°6.363 X 10-6 7.768 X 10-' 9.994 X 10-'

Sample III: 0.96250 gm, 99.999% purity, spheres 0.0965 ±0.0085 mm diameter

"799.40 105.95 1.359 X 10-'° 7.462 X 10-' 799.50 154.35 2.535 X 10-'° 9.287 X 10-' 717.00 73.65 2.132 X 10-" 3.218 X 10-' 799.50 799.40 201.50

245.95

2.846 X 10-'° 2.543 X 10-'° 7.428 X 10-' 7.968 X 10-' 788.80 854.50

73.00

73.35

5.686 X 10-"

1.011 X 10-'° 1.306 X 10-6 3.982 X 10-6 920.90 73.26 2.785 X lO-'o 1.056 X 10-5

" This datum was not used for the least squares equation or

plot since it was outside the limits of error using Chauvenet's

Criterion [15].

within the limits of error. The slope of the line is 0.04 ±0.07.

The phase boundary reaction constant as a function of temperature can be expressed:

SmoO., (99.9%), K= 7.4 exp (- 35.6 ± 0.5 kcal//?D

Sm.Oa (99.998%), exp (-31.3 ±0.8 kcaMRT).

The pressure dependence expressed as a slope can

be written as

samples of different size were used in order to check

the validity of the equation. It was assumed that if a surface process were important but not correctly

accounted for by the phase boundary reaction constant, different diffusion coefficients would have

been obtained. Since the agreement was good, both

sets of data were used to give the temperature de- pendence as shown in figure 7 and given in the

equation

Z)= 1.3 X 10-4 exp (-31.0±2.0 kcalIRT).

184
-1 0.0.

.

— 10.5.

NcI,Ot n—

X

.75 -I-

-8 0 .8 5 .9 0 1.05 lOOQ/T (°k)

Figure 7. Nd203, Temperature dependence ofD.

Again there appears to be no pressure dependence

as can be seen in figure 8 where the slope was cal- culated to be 0.08 ±0.06.

1.5 1.6 1.7 1.8 1.9 2.0 2.1

Figure 8. A^c^jO^, Pressure dependence ofD. The phase boundary reaction constants as a function of temperature can be expressed

Nd203(V),A: = 2.5x 102exp(-40.3±0.1 kcalIRT)

NdaOa (VI), K = 1 . 7 exp (- 3 1 . ± 2 . kcal//? T)

6.3. Er203 Experiment

It was found in an attempt to measure the dif- fusion coefficient of oxygen in EraOs that the rate of exchange was much slower than had been observed for Nd203 or SmaOs. While it was possible

to make an estimate of the diffusion coefficient, it was not possible to determine it accurately. The

design of the apparatus did not allow higher temperatures to be used in order to increase the measured rate of exchange. The estimated diffusion

coefficient at 1000 °C was lO-i^cm'^/sec.

7. Discussion

In the presentation of results the only errors

fisted were those obtained as deviations from a

line determined by a linear least squares treat- ment. The magnitude of systematic errors could

be calculated from the variation of the sphere size and inaccuracies in pressure, temperature, and

volume measurements, with the result that the dif- fusion coefficients could be in error by 10 to 15

percent, the heats of activation by ±4 kcal/mole,

and the logarithm of the pre-exponential factor

by ±1. The systematic errors would probably be

of the same order of magnitude as errors deter- mined from the scatter if different samples were

used for each experimental point.

The low values measured for the heat of activation and the lack of pressure dependence suggest

that extrinsic diffusion is important. It is possible

that the defect concentration is controlled by

impurities, such as Ca, Mg, and Si which have a maximum fisted concentration of 0.01 percent by

weigfit. Since tfie pre-exponential term in extrinsic

diffusion is controlled by the concentration of defects, the difference observed for Sm203 99.9 and

99.998 percent might just be due to the differences

in the concentration of these impurities.

If there was a difference in the defect concentration between the sfightly reduced EraOs and

spheres, it is possible to explain the different dif- fusion coefficients observed for these two samples

by a change in the pre-exponential term Do- For

example, if the value for the heat of activation for

slightly reduced Er203 [6] is assumed to be correct,

by using a pre-exponential term of 10~^ cm^sec,

similar to that measured for Sm2 03 and Nd2 03,

the diffusion coefficient is close to 10~'^cm^/sec

at 1000 °C as was observed for the exchange with

spheres.

Considering the heats of activation measured

and the above assumption, it is possible to explain

qualitatively the reported results. Since in the case

of extrinsic diffusion the heat of activation would

be the energy of motion, differences in the inter- atomic distances along the jump path would be

expected to produce differences in the activation

energies. For the same rare earth cation the density

increases or average interatomic distance decreases

as the structure changes from C to i? to ^-type

sesquioxide [16]. While the limits of error and tfie use of different cations makes it difficult to reacfi any definite conclusion, Nd203 (^-type) does fiave a greater activation energy than SmaOs (fi-type) as might be expected from the oxide type. While

no activation energy was measured by isotope ex- change for Er203, if the 47.8 kcal/mole [5] is correct, it must be assumed that the lanthanide

contraction has decreased the interatomic distances

to such an extent that even though the structure

is C-type, the energy of activation is greater than

that observed for the A-type and fi-type sesqui290-621 0-68— 13 185
oxides measured. This ionic radius effect can also be used to explain why the difference between

Sm^Os and NdaOs was not greater.

Before the diffusion experiments can be expected

to yield unambiguous information about the trans- port mechanism the effect of structure alone must

be explored. (Such a study is presently underway

with B and C-type GdOi..-,.) It will also be necessary

to clarify further whether or not the diffusion is indeed extrinsic and if so to what type of defect

it is due.

It is a pleasure to express appreciation to the

United States Atomic Energy Commission for their

support of this work. Dr. S. H. Lin provided many

interesting and enthusiastic discussions concerning the solution of the diffusion equation.

8. References

[1] L. Eyring and B. Holmberg, Adv. Chem. Ser.. 39, 46 (1963).

[2] B. G. Hyde, D. J. M. Bevan, and L. Eyring, Phil. Trans. Roy. Soc, London, Ser. A, 25 9 , 583 (1966).

[3] K. Vorres, Ph.D. Thesis, State University of Iowa, Iowa

City, Iowa (1958). K. S. Vorres and L. Eyring, "Rare

Earth Research" (E. V. Kleber, Editor) The MacmiDan

Co., New York (1961).

[4] U. Kuntz and L. Eyring, Kinetics of High Temperature

Processes, Ed. W. D. Kingery. Cambridge Technology

Press (1959), p. 59.

[5] J. Brett and L. Seigle, J. Less-Common Metals, 10, 22

(1966).

[6] C. D. Wirkus, M. F. Berard, and D. R. Wilder, J. Am.

Ceram. Soc, 501, 113 (1967).

[7] R. Haul, G. Diimbgen, and D. Just, J. Phys. Chem. Solid, 26, 1 (1965).

[8] P. Caro, Inorg. Chem.. In press.

[9] P. Jiru and J. Novakova, Coll. Czech. Chem. Commun.,

28, 1 (1963).

[10] P. Jiru, K. Kher, and J. Novakova, J. Catal., 2,479(1963).

[11] J. Crank, Mathematics of Diffusion, Clarendon Press, Oxford, 1956.

[12] P. C. Carman and R. A. Haul. Proc. Roy. Soc, A222, 109

(1954).

[13] R. Haul. G. Diimbgen. and D. Just, Z. Phys. Chem. NF.

31, 309 (1962).

[14] Handbook of Mathematical Functions — AMS55 (M. Abramowitz and I. A. Stegun, editors) U.S. Department of Commerce, National Bureau of Standards.

[15] A. G. Worthing and J. Geffner, Treatment of Experimental

Data, John Wiley & Sons, Inc., New York (1943).

[16] G. Brauer, Progress in Science and Technology of the Rare

Earths, Vol. 2, Ed. L. Eyring, Pergamon Press, New

York (1966).

186
Oxygen Transport During Oxidation

John B. Lightstone ' ' ^ and J. Paul Pemsler ^

A large volume fraction of potential short circuit paths for diffusion was shown to exist in the oxide

formed on Zr by oxidation between 500 and 1000 °C

in oxygen. The results of oxygen tracer experiments were also in agreement with a short circuit diffusion mechanism of oxygen transport. The as- sumption of line diffusion was shown to imply

intuitively reasonable values of both the grain

boundary thickness and the ratio of the oxygen

line diffusion coefficient to the lattice diffusion

coefficient. X-ray line broadening measurements were used

to estimate the grain size of the oxide. The as- sumption was made in the interpretation of the data

that strain contributions to the line breadth may be

neglected. The average grain was calculated to be

approximately 300 A. The small grain size and a possible high dislo- cation density associated with a martensitic trans- formation imply a high volume fraction of potential

short circuit paths for diffusion. The mechanism of

oxidation was studied in an attempt to demonstrate

that transport of oxygen along these short circuit paths is the predominant mechanism of film growth. Zr was oxidized at 500, 600, 750, and 910 °C in 0|^, and the oxidation continued in OJ**. A new

method of general application in diffusion problems was used for studying the movement of the oxygen

tracer. The oxide was sputtered in a beam of argon

ions at a rate of 1 micron per hour into a mass

spectrometer. The O'^ profile was determined with

a resolution of 0.015 microns.

The fraction O^** exchanged in a volume very near the surface was always less than unity. The

fraction increased with time in Ol** for the same

initial time in at a given temperature.

Arguments are presented against an explanation

of these observations in terms of a surface ex- change limited reaction, or in terms of an O"* interstitial diffusion followed by exchange with O'** in normal sites. However, the evidence against a

' Ledgemonl Laboratory, Kennecott Copper Corp., Lexington. Mass. 02173. ^Now at Union Carbide Corp., Linde Division. Tonawanda, N.Y. 14152.

surface exchange fimited reaction is not conclusive.

The observations of the fraction O'** exchanged

in a volume very near the surface are explained by

suggesting that many of the oxygen atoms are not taking part directly in the diffusion process. Transport is occurring predominantly through easy dif- fusion paths. There is also a slow exchange between

O'^ in these short circuit paths and O'^ in the

lattice, so that the fraction exchanged near the

surface increases with time.

The fraction exchanged after very short times

such that there is a negligible contribution from

lattice diffusion to the diffusion profile is used to estimate the volume fraction of short circuit paths.

An estimate of the volume fraction of short circuit paths obtained in this way is combined with a value

of the average grain size from x-ray fine broadening

measurements to derive an apparent grain boundary

thickness on the assumption that grain boundaries

are the only short circuit paths available. At 600 °C

this implies a grain boundary thickness of 14 A. This result is similar to values reported in metals,

and is intuitively reasonable. This is regarded as evidence in favor of the model.

Observations of the O'** profile after very short

times in O^** were used to derive a fine diffusion

coefficient for oxygen. The increase with time of

the fraction exchanged in a volume near the surface was used to calculate an oxygen lattice diffusion

coefficient. The ratio of the fine diffusion coeffi- cient to the lattice diffusion coefficient was 9 X 10'' at 600 °C. This ratio is similar to values obtained

in metals, and is cited as indirect justification of

the model.

Nickel oxidation at 1000 °C was studied by a similar tracer technique. The fraction exchanged in a volume very near the surface was in this case unity.

The shape of the diffusion profile was consistent

with a model of material transport in which there

is a significant contribution from oxygen lattice diffusion.

187

Mechanical and Dielectric Relaxation of Hopping Electrons and the Ionic

Defects in Reduced Rutile (TiO^_^)

W. W. Scott and R. K. MacCrone '

School of Metallurgy and Materials Science

University of Pennsylvania, Philadelphia, Pennsylvania 19104

Non-stoichiometry in transition metal oxides involves the formation of ionic defects. For oxygen

deficient n-type oxides, the primary ionic defects are oxygen vacancies and/or cation interstitials. These ionic defects are positively charged and at lov^^ temperatures trap electrons into states centred on one of several nearest neighbour cation sites. Some general features of the electron localization are briefly discussed. An electric field or a mechanical stress may induce electron hopping betv^feen two such cation

sites and give rise to a discrete dielectric or mechanical relaxation. The orientation dependence and

"selection rules" of such relaxations provides a powerful method of investigating the ionic defects and

complexes. In the particular case of reduced Rutile several dielectric relaxations and their corresponding anelastic relaxations are observed. The orientation dependence, and impurity dependence, of the

dielectric relaxations indicates the presence of oxygen-vacancies in association with extrinsic trivalent impurities and also Ti + ^ interstitials which may be associated with two impurities.

y

Key Words: Cation interstitial, nonstoichiometry, oxygen vacancy, rutile.

1. Introduction

Mass transport in oxides occurs by the diffusion of ionic defects. In all but the simplest cases, more

than one kind of ionic defect is involved, since at high temperatures various interactions may occur and defect complexes may be formed. Thus, ex- periments which measure the combined effect of several defects cannot easily yield unambiguous

conclusions about the individual defect types in- volved. Selective experiments which enable specific

defect types to be identified and their concentrations determined are therefore of great importance.

Dielectric and anelastic relaxation measurements

near and above room temperature have proved to be valuable in identifying specific ionic defects,

where the ion itself moves between equilibrium

positions. This paper describes how at low temperatures in a reduced transition metal oxide,

TiOa, electron hopping gives rise to dielectric and anelastic relaxations, and how these relaxations

may be used to characterize the ionic defect types.

The experimental method should be applicable

to other transition metal oxides.

The basic mechanism involves the ionic defect centres (e.g., oxygen vacancies) which are positively charged and which trap electrons at low

temperatures. These electrons are not in hydrogenic

states centered about the particular ionic defect

but in states centered around one of several sur- rounding nearest neighbor cations. This can be

schematically represented as follows:

An equivalent configuration of equal energy but

of opposite dipole moment is

An alternating electric field will induce backward

and forward electron hopping, which under ap- propriate conditions of temperature, frequency,

and orientation will give rise to a Debye type

dielectric relaxation [1-5]. ^ A theoretical dis- cussion has been given by Frohlich, Machlup

and Mitra [6J.

2. Experimental Results and

Discussion

Dielectric relaxation measurements at low temperatures with the electric field parallel to [001] have

shown [7, 8] the existence of two pronounced re-

' Present address: Department of Materials Engineering, Rensselaer Polytechnic

Institute, Troy, New York 12181. Figures in brackets indicate the literature references at the end of this paper.

189
10 _ SPECIMEN P ( STOICHIOMETRIC) q——"

- 5io-0-0—O o O °^ 10 kHz

i I I I I I . 5 10 15 20 25 30 TEMPERATURE, °K

Figure 1(a). Typical dielectric loss as a function of temperature of reduced Rutile, showing two relaxation processes. E 1 1 [001].

^= 1.0±0.5 X 10-3 gY ^jti^ pre-exponential

factor T* = 10"** sec. The higher temperature peak. Peak II at 17 °K,

has been shown to be most hkely due to electron

hop I at the Ti+3 interstitial, which may be asso- ciated with two trivalent substitutional impurities

as indicated in figure 3. The relaxation time for

this peak was found to have an activation energy ^ = 2.0±0.5 X 10-2 gY ^jjI, pre-exponential

factor T*= 10-^-10 -« sec. The identification of the two relaxations shown in figure 1(a) with the specific electron hops and ion cores indicated was deduced from experiments

in which the orientation of the electric field E was

varied with respect to the crystallographic axes and from experiments in which the trivalent substitutional impurity was varied.

Figure 1(b) shows the behaviour when the eleclaxation peaks, figure 1(a) (The temperature region

above 30 °K is obscured in these measurements

due to the formation of a Shottky barrier layer as the Fermi level rises with temperature.). A comprehensive discussion of the structural origin of these relaxations has been given [7, 8]. The salient features are summarized below; the details will be

found in refs. 7 and 8. The low temperature peak. Peak I at 6 °K, has

been shown to be due most likely to electron hopping at an oxygen vacancy in association with one

trivalent substitutional impurity. This complex

defect has two, almost equi-energetic ion core

configurations. These core configurations are shown

in figures 2(a) and (b), which also indicate the

proposed electron hops. The relaxation time t for

this peak was found to have an activation energy

lo-'Lv—J 1 1 ^ ^ 1 0 5 10 15 20 25 30 TEMPERATURE, °K

Figure 1(c). Dielectric loss vs. temperature for specimen con- taining increased trivalent substitutional impurity. E||[001].

190
trie field is parallel to [100], the specimen with

electric field parallel to [110] behaving similarly. In both cases the higher temperature peak is absent

since the hopping mode 2, figure 3, cannot be

excited for these field directions.

Figure 1(c) shows the behaviour of a specimen

in which the trivalent substitutional impurity had

been increased. In this case Peak I is absent due

to the proposed formation of (Ov — Y-z) neutral

complex defect centres [9] which cannot trap electrons. The formation of this (neutral) complex

defect probably follows the equations

0?2 + Y-i/^,(r)(0,-Y)+i

(O,- y).+' + Y- 'K-ATnOr - Y.9).

By increasing the concentration of the effectively

negative trivalent impurity the equilibrium

is driven towards the right, since Haul and Dumbgen [10] have shown experimentally that the concentration of oxygen vacancies remains fixed.

Trivalent Impurity

ELECTRON HOP

|/^[hM]l

i

[lOO] 2.230 [no] 3 248

[OOl] 1.480

0 Trivolent extrinsic impurity Q Oxygen voconcy

O Normal lottice titonium ion ® Titonium interslitiol

ELECTRON HOP ,A

e

[lOO] 0

2 [no] 0

[OOl] 2 959

Figure 2. Oxygen vacancy with one associated trivalent impurity.

(a) [100]. [110|. and [001) components of dipole moment, fi (b) only [001| component o{ fji.

|/^[t,Mj|

ELECTRON HOP e

'

[lOO] 0

1 [llO] 0

[OOl] 2,959

[lOO] 4594

^ [llO] 3.248

[OOl] 0

Figure 3. Titanium interstitial associated with two trivalent impurities.

The intermediate temperature peak. Peak III (which is obscured in figure 1(a)) is also due to electron hopping at some ionic core. Since this relaxation shows the same orientation dependence

as Peak I, and apparently does not depend on the

trivalent impurity concentration, this ion core has

been proposed as a simple substitutional pentavalent impurity present in our boules. The relaxation time for this peak was found to have an

activation energy = 8.0 ± 0.5 X 10"^ eV and t*

10-7 gg(. There are two major theoretical aspects that are

of interest concerning the localization of the elec- trons and the intercationic hopping as envisaged:

(i) the symmetry properties of the defects

and the selection rules for observable

relaxations and

(ii) the nature of the electron states.

2.1. The Symmetry Properties

The symmetry properties and selection rules of

defects in general, have been discussed by Wachtman [11], and by Nowick and Heller [12]. These

references should be consulted for further details.

2.2. The Electron States

Consider the oxygen vacancy shown in figure 2(a). If the cation sites were strictly equivalent, the

ground state wave functions of the electron would

be symmetric or anti-symmetric combinations of wave functions centered on the cation sites. Transitions between these non-polar states would thus

not give rise to a dielectric relaxation. As discussed

in detail in [8], transitions between the ground state and some polar excited state will give rise to a relaxation but cannot be the cause of the dielectric relaxations observed here; the transition rate would

be too rapid. Therefore, the supposition that the

two cation sites are equivalent must be abandoned.

The non-equivalence of the two cation sites may

arise in one of two ways:

(i) the "reaction field" of the lattice,

polarized by the electric dipole, is

191
different at the two cation sites [8], or

(ii) the relative displacements of the two

cations with respect to some symmetry

centre are different.

It is not known which of these two effects is dominant. In any event, a perturbation on the non- equivalence of the cation sites may be produced

by the application of a mechanical stress in an

analogous way to the perturbation produced by

an electric field. To specifically check this prediction, low temperature mechanical loss measurements were made on a specimen of lightly reduced rutile in an

apparatus described elsewhere. The specimen

orientation was such that an alternating compressive and tensile stress was applied along the [101]

direction. This direction was chosen since elec- trons at all three of the proposed defect centres

will be active if the centres are present. The re- sults are shown in figure 4(a). (Note that in these

TEMPERATURE, "K

Figure 4(a). Mechanical loss versus temperature in reduced

rutile {dotted line ) showing two relaxation processes.

The correlation with the dielectric relaxations (solid line) should be noted. specimens Peak I is not present due to the presumed

large trivalent impurity content [8].) The one to one correspondence between the dielectric relaxations and the mechanical relaxations is most

satisfactory. It is interesting to note that, where

comparison is possible i.e.. Peak II and Peak III, the activation energies Q and pre-exponential

factors T* are the same for the mechanical and

dielectric relaxations. The values of Q and t * de- scribing electron hopping in Rutile are comparable

to those reported by Volger [5J for electron hopping

in Fe-iOa. The mechanical loss measurements reveal in addition other relaxations at temperatures 42 and

55 °K (fig. 4(b)). These relaxations (provided of course that they are electrically active) have not been observed in vacuum reduced rutile since the

formation of a Schottky surface barrier obscures

_ 20

OO 15

-

/ ^\,l6IOHz

/ \^ SPECIMEN W

/ (LIGHTLY REDUCEDl ~

/ ^— [lOll / 1720 Hz

z

\

, // SPECIMEN X

( VERY LIGHTLY REDUCED) [lOi]

1 1 1 1 0 20 40 60 80 100 TEMPERATURE,°K Figure 4(b). Mechanical loss vs. temperature in reduced rutile for two states of reduction. bulk dielectric phenomena at temperatures above

30 °K. Mechanical measurements do not have this disadvantage.

The most powerful approach in identifying any

active defect is to study the symmetry properties

of the resuhing relaxations. Such an investigation

is currently in progress on the peaks at 50 °K. Any

serious ion core model must await the results of these experiments.

3. Conclusion

Low temperature dielectric and mechanical re laxation due to intercationic hopping at ionic defect centres has been shown to be a useful method of

investigating ionic defect centres in transition metal

oxides. Using this technique, the orientation de- pendence of the dielectric relaxations that occur

in reduced Rutile indicates the presence of oxy- gen vacancies in association with one trivalent substitutional impurity (fig. 2(a) and (b)) as well

as the presence of Ti+^ interstitials (fig. 3). A one

to one correlation between dielectric relaxations

and mechanical relaxations has been established

for two distinct electron hops. In addition, new

electron hops at ~ 50 °K have been found.

From a theoretical point of view, the localiza- tion of electrons in this way requires that two cation

sites, equivalent in the lattice, be unequivalent

when adjacent to an intrinsic ionic defect with an

electron locahzed on one of the cations. In this connection the mechanical relaxation indicates

that a mechanical stress, as expected, alters the

potential inequivalence sufficiently to induce

transitions between the two sites. It is not known

whether the cation site inequivalence is due to polaron effects or to electrostatic effects within

the ion core.

Finally, the magnitude of the dielectric relaxations may be used to determine the number of

electrons involved quite accurately (on the order

of 10'** cm"^ in fig. 1), since the electric dipole

192
moments can be reasonably estimated. There is probably a fairly good correlation between the

electron concentration and ionic defect centre concentration, although at these temperatures the

effect of other polar or non-polar donor or acceptor

states may be important in the Fermi statistics of

these wide gap materials.

The authors would like to express their appreciation to Dr. Jerome R. Long for helpful discussion

and the editors of The Physical Review for per- mission to reproduce figures 1 to 3. The financial

support of the Advanced Research Projects Agency

is also gratefully acknowledged.

4. References

[1] J. W. Verwey, Semi-conducting Materials, Butterworths,

p. 151 (1951).

[2] J. Volger, J. M. Stevels, C. van Amerongen, Philips Research Reports, 10, 260 (1955).

[3] N. R. Freyman, et al., J. Phys. et Raduim 17, 806 (1956).

[4] W. A. Joffe, Silikattechnik 10, 599 (1959); Fiz. Tv. Tela

4, 669 (1962).

[5] J. Volger in Progress in Semiconductors, Vol. 4, Ed.: A. Gibson, Wiley, p. 205 (1960).

[6] H. Frohlich, S. Machlup and T. K. Mitra, Phys. Kondensierten Materie 1,359(1963).

[7] L. A. K. Dominik and R. K. MacCrone, Phys. Rev. 156,

No. 3, 910 (1967).

[8] L. A. K. Dominik and R. K. MacCrone, Phys. Rev. 163, No.

3, 756 (1967).

[9] * The oxygen vacancy Of^ is effectively positive with a double charge.

A trivalent substitutional impurity is effectively negatively

charged in the rutile lattice, and is given the symbol

[10] R. Haul and G. Dumbgen, J. Phys. Chem. Solids 26, 1 (1965).

[11] J. B. Wachtman, Jr., Phys. Rev. 131, 517 (1963).

[12] A. S. Nowick and W. Heller, Advances in Physics XIV,

No. 54, 101 (1964).

193

Purity and Perfection of Research Specimens of Oxides ^

J. W. Cleland

Solid State Division, Oak Ridge National Laboratory

Oak Ridge, Tennessee 37830

Specific examples of the effects of pre-purification, crystal growth, characterization, and research

involving oxides in the Research Materials Program of the Solid State Division of the Oak Ridge

National Laboratory were selected to indicate the current status of the purity and perfection of these

particular materials, and to examine the primary experimental problems that stand in the way of

better research materials. The coded and filmed collection of data sheets, abstracts, reports, open

hterature publications, and bibliographies of the Research Materials Information Center, various

publications of the Electronic Properties and Defense Ceramics Information Centers, and the contents

of a number of journals and books gave information on the initial purification, crystal growth, and

final characterization of some general classes of oxides, as well as a number of specific materials of

interest. A number of individual research groups were contacted to obtain current information on the

purity and perfection of those oxides of particular interest to their own research program.

Key Words: Characterization, crysteJ growth, oxide crystals, purity.

1. Introduction

The purpose of this survey was to ascertain the

present status regarding the purity and perfection

of research specimens of oxides. Progress in un- derstanding the fundamental properties and the

ultimate range of properties in many materials has

been severely Umited by a lack of pedigreed

samples of sufficiently high purity and crystalline

perfection. Extensive tests on ultra-pure specimens, or on specimens that contain a known amount

of a specific impurity, are required.

The Research Materials Program of the Sohd

State Division of the Oak Ridge National Laboratory (ORNL) is concerned with developing the

necessary techniques required for the growth of

ultra-high purity and of controlled impurity re- search specimens (usually single crystal) of a composition, quality, purity, and perfection required

in materials research. The choice of materials for consideration has been made, for the most part, on the basis of the potential gain to a fundamental

understanding of a given material that may result from an intensive investigation by a number of

research groups, and on the basis of an expressed

need for a specific research material. Certain re- search materials that have not been investigated

previously are being produced with new techniques

of purification and crystal growth.

The entire program involves a number of research

divisions at ORNL and other AEC installations.

' Research sponsored by the U.S. Atomic Energy Commission under contract with Union Carbide Corporation.

The Analytical Chemistry Division has provided

a variety of special analytical techniques, and a program of research on improvement in analytical

techniques. Most of the work on crystal growth has

been carried out by research groups within the

Metals and Ceramics and the SoUd State Divisions

of ORNL, and these groups have also conducted

a wide variety of physical property measurements

that characterize the end-product. Selected samples have been supplied to other research groups

within and without ORNL for comparative analyses,

cooperative measurements that increase the characterization, and for special experiments of

interest to the AEC.

Information concerning the availability of and

need for high-purity research materials has always

been difficult to obtain; hence, the Research Materials Information Center (RMIC) [Ij- was established as a part of the Research Materials Program

to collect and provide information on the initial purification, crystal growth, final characterization,

and availability of research quality specimens to both producers and users. Ready access to its ac- curate up-to-date listing of available materials

eliminates much duplication of effort to produce

materials that are already available elsewhere. A

simultaneous fisting of research materials that are desired but are not available has served to focus

the attention of researchers on promising new

areas of investigation.

The American Ceramic Society has published

a compilation [2J of selected engineering property

^Figures in brackets indicate the literature references at the end of this paper.

195
data of the physical, thermal, and mechanical

properties of a large number of single and mixed

oxides. The separate hsting of 1026 data sources

or references is of value in obtaining information on purity and perfection that was not a part of the

collection of engineering properties. The Defense

Ceramic Information Center (DCIC), Battelle Memorial Institute, Columbus, Ohio, collects,

interprets, and disseminates technical information

to qualified requestors on materials composition,

processing, fabrication, property and performance

data, and fundamental aspects of processing and

behavior. The Electronic Properties Information

Center (EPIC), Air Force Systems Command,

Hughes Aircraft Company, Culver City, California,

has compiled information on the electronic properties of ALO3, BeO, MgO, CdO, PbO, and ZnO

in separate data sheet publications that are avail- able from the Defense Documentation Center,

Alexandria, Virginia.

Review articles [3| by E. A. D. White and R. G. Bautista and J. L. Margrave were of considerable

value in obtaining information. A report [4| of the

Materials Advisory Board, the International Conference on Crystal Growth [5], and the newly formed

Materials Research Bulletin, Journal of Crystal

Growth, and Journal of Materials Science provided

additional information of value. The coded and microfilmed collection of data

sheets, abstracts, reports, and open literature publications of the RMIC, and the data sheets and

data sources of the other publications named above

gave information on the initial purification, crystal

growth, and final characterization of general classes

of oxides, as well as specific materials of interest.

It is evident that individual authors, sponsoring

organizations, and journal referees or editors must

share the blame for minimizing or deleting the in- formation that concerns the actual purity or per- fection of the specific material under investigation.

Most of the publications referred to the material

as good quality, reasonable purity, or as single

crystal; however, the actual sample was seldom

identified in quantitative terms of specific impurity

content or lattice perfection. It is hoped that the

above mentioned new publications (and the old)

will alleviate this problem.

A number of individuals and research groups were contacted to obtain up-to-date information on the purity and perfection of those oxides of

particular interest to them. No attempt was made

to contact every individual or group, hence, the

information as received was incomplete; however,

the cooperation of those who furnished information

contributed in large measure to the aim and intent of this survey.

The RMIC, already described, answered more

than 2400 inquiries from 940 individuals in 27

countries during the period from July. 1966, to July, 1967. Table 1 is a tabulation of requests re- ceived for information regarding the availability

of single crystals of a particular material, arranged

in order of the number of requests during the period

from July, 1966, to July, 1967. This hst does not include the requests for other information, such

as purification, crystal growth techniques, or

specific property measurements, but only the direct

requests for availability of single crystals.

2. Specific Materials

2.1. Magnesium Oxide

Single crystals of MgO are advertised as avail- able from Electronic Space Products, Inc., General Electric Co., Norton Co., Monocrystals, Inc., Muscle Shoals Electrochemical Corp., SemiElements, W. & C. Spicer. Ltd., England, and yet

other sources of a commercial or research nature. A large number of detailed analyses of both start- ing material and final product appear in the litera- ture; however, the end result is 99.9 percent

(3/V) or at most 99.99 percent (4A') purity. This is not surprising since most crystal growers have

employed the same type of starting material, and

since a few companies have served as producers

for a number of crystal brokers or vendors.

An example of the complexity of analytical

techniques for materials is that of a long-term contract between the Jet Propulsion Laboratory and

the Battelle Memorial Institute for The Developrnent of Chemical Analysis Techniques for AdTable 1. Single Crystal Requests Through The Research

Materials Information Center (July 1966-July 1967)

Crystal

Number of Requests

Crystal

Number of Requests

MgO 17 KCl 8 ZnO 16 Te 8 BaTiOa 15 Fe304 7 CdS 14 CaF. 7 Ge 12 Ni 7 graphite 12 NiF. 7 ZnS 12 Si 7 PbTe 11 SrO 7 Ti02 11 FeF.> 6 ZnSe 11 MnO 6 ZnTe 11 W 6 BaO 11 AIN 5 HgS 10 PbS 5 NiO 10 PbSe 5 BaS 9 SrTe 5 ThOs 9

196
vanced Materials. The intent of this contract is the complete chemical characterization of the

oxides of Mg, vVl, Ca, and Zr, and the carbides of

Ta, Hf, Nb, and Zr. The progress to date on one

starting material (Fischer M-300 magnesia powder)

is that it has been standardized in terms of F,

Si, S, CI, and Ca. The primary technique for single crystal growth

of pure and of doped crystals of MgO is the arc

fusion between carbon electrodes of precalcined

or sintered material at temperatures in excess of

2800 °C. A. Rabenau [6] has described a typical

furnace, melting procedure, and technique for selected doping. The Atomic Energy Research

Agency, Harwell, England, had a contract with

W. & C. Spicer, Ltd., for arc fusion growth of

single crystals of MgO, and the Research Materials Program at ORNL has purchased the right

from that company to duplicate the furnace design

and crystal-growth techniques.

Table 2 indicates the analyses of magnesia

powder as obtained after isostatic compaction

and following arc fusion. It is evident that the

initial purities of the starting material and the

electrodes, and the maintaining of purity during

calcining, sintering, isostatic compaction, fusion,

ball mill grinding, and refusion were the factors

that determined the purity of the end product.

The MgO crystals were of reasonable purity

(4A'^) and perfection; however, many of the melts

gave crystals that exhibited an internal cloud or murkiness that may have been due to hydroxyl or water vapor in the starting powder. The Research

Materials Program at ORNL will investigate [7] some additional prepurification techniques, pro- longed vacuum baking at an elevated temperature

before isostatic compaction, a close control of

purity during compaction and fusion, and high temperature vacuum annealing of the fused crystals

in an attempt to improve the quality.

Figure 1 indicates the optical absorption coeffi- cient at 2850 A of MgO quenched in argon and irradi- ated with ^"Co gamma rays versus the concentration of iron (from chemical analysis) for a wide

variety of single crystals measured by Y. Chen and

W. Sibley of the Solid State Division of ORNL [8]

.

Also shown are similar data for MgO with 200 to 5000 ppm of iron, obtained by R. W. Davidge of

AERE, Harwell, England, after heating the crystals

in air at 1400 °C [9]. The figure shows that the optical absorption coefficient depends markedly on

the concentration of iron in the MgO.

It must be stated that most of the published work

on diffusion, plastic deformation, and optical ab- sorption studies on MgO would have benefited from

better crystals. The status-quo is that of single

crystals of several tens of cubic centimeters max- imum size, 4A'^ purity, internal cloud due to hydroxyl

or water vapor, and lots of experimental problems

to solve before better crystals can be provided.

2.2 Zinc Oxide

The final report of a contract between the Advanced Research Projects Agency and the Radio

Corporation of America on "The Synthesis and

Characterization of Electronically Active Materials" contains the statement by L. R. Weisberg

that "ZnO was chosen for study because it is versatile, useful, and enigmatic". The author continues that "The means and mechanisms of energy

transfer in ZnO are still obscure". The files of

the RMIC contain many references to crystal growth

of ZnO during the past five years by hydrothermal,

flux growth, traveling solvent, vapor reaction,

thin-film sputtering, and other techniques. The

Table 2. Analysis of magnesia powder and single crystals of MgO (ppm) compact and crystal analyses at ORNL

Impurity Ca Na Pb Ba Mn Zr As Cr

Powder ^ 35

29 20-30

25 <0.8

<0.5

40 < 20 <23

1.5 < 1

trace < 1 1-510 <3

< 5 Compact

Single crystal '

25

1-3 24 < 5

Impurity Ti N P S Fe Si Al Zn

Powder ^ 25

8.5 < 3

250

140 < 2

8

11 2-10

20

76 15-39

80

50 30-40

20 <50

1-10 Compact

Single crysted <30<5

< 1 1-5

^Powder and analysis from the Kanto Chemical Co., Japan. ''Same powder, compacted at Atomic Energy Research Agency, Harwell. England; analyzed at Oak Ridge National Laboratory. •^Average values for crystals from 18 arc fusions by W. & C. Spicer, Ltd., England; analyses at Oak Ridge National Laboratory.

197
sooo

20 oo

Eu 500 —

10 '2 5 10'^ 2 5 10"^ 2 5 10"^ 2 £ CONCENTRATION OF Fc FROM ANALYSIS (Nfj/N^j)

Figure 1. Variation of absorption in MgO at 2850 A wavelength with iron content. O— saturation level on y irradiation (Chen and Sibley); C —saturation level on heating in air at 1400 **C.

lO

Airtron Division of Litton Industries has produced

ZnO hydrothermally, and this material is commercially available. Y. S. Park and D. C. Reynolds

[10] have grown crystals by subliming both ZnS

and ZnSe to have them react with oxygen in an

argon gas flow. Figure 2 is a photograph of single crystals of

ZnO, available commercially in research quantities [11 J; Table 3 gives the results of an emission

spectrograph analysis (typical data but not guaranteed values). The status-quo for ZnO, therefore,

is that of raw, uncut, hexagonal crystals up to 6 mm diameter, with less than 50 ppm impurity

content as detected by emission spectroscopy or available in the Z-axis wafer form, either undoped

or doped with lithium. It is hoped that a commercial

availability of research specimens of reasonable

purity will serve to make ZnO less enigmatic.

2,3,fialcium Oxide

Single crystals of CaO of 2N or 3N purity are

advertised for sale from Electronic Space Products,

Inc., Muscle Shoals Electrochemical Corp., SemiElements, W. & C. Spicer, Ltd., and yet other

commercial producers. Table 4 represents typical

data for the Spicer product. These crystals are

listed as 3N purity, and they vary from colorless transparent to cloudy or milky. It has been stated

that the cloud can be removed by annealing at 1800 °C. F. Galtier and R. CoUongues [12] have

recently reported the production of single crystals

of CaO by the Verneuil method in a plasma furnace.

They state that the crystals are as large as 7 mm

diameter by 20 mm long, are perfectly transparent,

and that they exhibit a remarkable chemical inertia

in that they can be stored under water for several

weeks without change. No information is available

as to the actual purity and perfection of these crystals; however, the fact that they are transparent

and unaffected by water would suggest that a considerable improvement has been made in the quality

of single crystals of CaO.

198
Figure 2. Single crystals of ZnO.

2.4 Barium or Strontium Oxide

During the period from July, 1966, to July, 1967,

the RMIC received 12 requests for single crystals

of BaO; however, they had no information concerning a commercial or research producer. R. J. Gambino [13] produced single crystals of Eu-doped

BaO or SrO by a modified Verneuil technique,

and G. E. Pynchon and E. F. Sieckmann [14] grew

single crystals of "good optical quality" in a graphite

resistance furnace by the method of ohmic heating;

however, the maximum size of these crystals was

listed as 1 mm'^. The pioneering effort of the group

at Cornell University [15] has not been continued

due to a lack of research quality specimens.

2.5 Zirconia and Hafnia

Analytical data exist for single crystals of zir- conia, Zr02, and hafnia, Hf02, that have been

produced in a number of modifications. Pure

zirconia has been used to produce crystal fragments

up to 7 cm^ volume with no visible crystal planes

[16] as: (a) opaque green to white crystals, (6) gray

to black crystals that possibly contain hydrogen

in the lattice, and (c) gray to black crystals with a surface coating of golden appearance [17]. On

calcination both (6) and (c) revert to (a). Transparent crystals up to 0.5 cm^, varying in color from

pale straw to dark brown, have been grown [16] with additives of 9 percent wt. calcia or 14.5 per- cent wt. yttria. The maximum impurity level is about 700 ppm excluding hafnia, and Ruh and

Garrett [18] have published a full analysis for about

90 elements. Chase and Osmer [19] have produced

single crystals of zirconia and hafnia from a PhF^

fused salt system, and have also reported on flux crystal growth of other materials [20].

2.6 Titanium Dioxide or Oxide

The files of the RMIC hst 9 commercial and 7 research producers of single crystals of material

that are identified as Ti02, Ti203, or TiO and are labeled as reasonably pure or of good quality. The

literature as available, however, is almost devoid

of any indication as to the purity and perfection

of such crystals; not even the stoichiometry is indicated. Additional work involving the Verneuil,

flame fusion, float zone, hollow cathode plasma

torch, or chemical transport [21] techniques are required before one can discuss the purity and

perfection of these materials.

199
Table 3. Emission spectograph analysis of single crystals of

zinc oxide (ppm) made by the 3M Co., Dielectric Materials and

Systems Division

Table 4. Analysis of single crystals of calcium oxide (ppm) from W. & C. Spicer, Ltd.

Element

Cu.

Cd.

Ti.. v.. Cs.

Sn. Mo

Be.

Al..

Bi..

In.. Ge.

Fe.

Cr.

Ni.,

Si.. Sb. Mg

Pd. Mn

B... Te. As.

Ba.

Li.. K.. Na.

Detection

limit

< 1.0 < 1.0 3 < 1.0 < 1.0 < 1.0 < 1.0 < 1.0 < 1.0 3 < 1.0 3 < 1.0 < 1.0 < 1.0 < 1.0 510 < 1.0 3 < 1.0 < 1.0 10

10

3 < 1.0 3 < 1.0

Quantity

< 1.0

< 1 < 11.5

<3

< 1 < 1 < 5 < 1

< 13

3

(Typical data but not guaranteed values. The sample was ground in a boron carbide mortar and pestle.)

3. Crystal Growth Techniques

3.1 Arc-transfer Method for Metal Oxides

A major advance in crystal growth of oxides has

been the arc-transfer method developed by J. R. Drabble and A. W. Palmer [22] of the University

of Exeter, England. They strike a d.c. arc between

a sintered feed rod (cathode) and the crystal being

grown (anode); the oxides of Co, Fe, Ni, Ti, and V

have been grown. B. A. Smith and I. G. Austin

[23] of the University of Sheffield, England, have

extended the arc-transfer technique to produce pure

and Li-doped crystals of NiO and CoO. Figure 3

is a photograph of typical crystals of NiO and CoO,

and tracking in NiO. The typical crystal is 5-9 mm

diameter X 1-2 cm long.

Several other systems were explored using the

arc-transfer method. Transport was usually ob- served in experiments involving Fe, Mn, Ti, V, Cr,

and Cu but the product often contained mixed

phases of more than one oxide. Only negligible

Impurity Al Fe Mg Mn Si Sr

< 50 < 200 <50 < 100 < 100

Ag Ba Co Ce Cd Be

Impurity Be Cu Pd Ni Na

Cr Ga Ti Pb Cr Mo Sn

<50 < 100 < 100

Figure 3. Metal oxide crystals prepared by the arc-transfer method.

(a) NiO; (b) CoO; (c) Tracking in NiO; (d) Cleaved boule of CoO (1 cm/div).

growth was observed for MnO, Cr203, AI2O3, and

Zr02. There is little if any data available on the

purity and perfection of the oxide crystals pro- duced by arc-transfer.

3.2 Epitaxial Deposition of Metal Oxides

Also of interest is the work of Robinson, White,

and Roy [24] of the Materials Research Laboratory

at the Pennsylvania State University on epitaxial

deposition of NiO, CoO, FeO, C03O4, and Fe304

200
on to single crystal MgO by the hydrolysis of the

gaseous anhydrous metal bromide. Repeated

deposition has permitted the growth of crystals

up to 1 mm thick, and with a substantial increase

in purity. The dislocation count for NiO was only

103 to 105 cm -2.

3.3 Internal Centrifugal Zone Growth of

UO2 and 50-50 ThOo-UOa

A second major advance in the growth of single

crystal oxides is the internal centrifugal zone growth

technique developed by A. T. Chapman and G. W.

Clark [25] of the Research Materials Program at the Oak Ridge National Laboratory. Figure 4 is a schematic drawing of the equipment to grow single

crystals of UO2. One starts with a sintered rod

having a void in the bottom to permit insertion of

a (single crystal) seed. The thin metal shields of Mo

are used for inductive coupling to the coil until the temperature is such that the coil couples di- rectly to the material. The dual action of rotation

and translation circulates the molten material

inside its own crucible, and deposits a single crystal

with the preferred orientation of the seed crystal.

This technique should be applicable to any high

melting point material whose electrical conductivity

at elevated temperatures will permit eddy current

heating. Large grains of 50-50 Th02-U02 have

been grown, also single crystals of isotopically

Figure 4. Equipment for the internal centrifugal zone growth

of single crystals of uranium dioxide.

enriched UO2 for which economy of material is enhanced by retention within its own crucible.

Figure 5 is a photograph of a cross section through

a typical ingot.

There are 6 commercial and 9 research pro- ducers of single crystals of UO2 listed in the files of the RMIC, and A. T. Chapman [26] is actively en- gaged in a program involving a comparison of material from different sources by etch pit, electron

diffraction, and electron microscope techniques.

This work was discussed at a recent meeting of the

Nuclear Society of the American Ceramic Society

and will be submitted for publication.

3.4. Flux Growth of ThOa and CeOa

A major interest of the Solid State Division at the

Oak Ridge National Laboratory is electron spin resonance (ESR) measurements on single crystals

of Th02 and Ce02 as doped with Gd, Dy, Er, or Yb.

Figure 6 is a photograph of a single crystal of ThO^,

an optically transparent tetrahedron about 3 mm

on edge, grown from a molten lithium ditungstate

solvent by C. Finch and G. W. Clark [27] of the

Research Materials Program at ORNL. These

crystals contain less than 500-ppm W and less than

50-ppm Li as major impurities. Lanthanide- or uranium-doped single crystals of Th02 and Ce02

have also been grown, and the glove box facilities of the Transuranic Research Laboratory at ORNL

are currently being used to incorporate Cm, Am,

and other transuranic elements. R. C. Linares [28]

Figure 5. Cross section through an ingot of uranium dioxide grown by internal centrifugal zone growth.

290-621 0-68— 14 201
0 12 3

mm

Figure 6. Octahedral ThO > single crystal grown from

Li 2O 2WO3 solvent at 1280 °C.

reported on the growth and properties of ThOi and

Ce02 as possible laser materials. Three different

flux systems were used, and the optical radiation

damage and possible laser properties were ex- amined. The author concluded that a laser crystal can probably be produced if one can devise a technique to remove excess lead from the system.

3.5. Electron-Beam Float-Zone Growth of

Laser Host Materials

Sapphire (alpha alumina) and ruby (Al203:Cr^+)

have been produced by flame fusion, hydrothermal,

flux, Czochralski, electron-beam float-zone, hollow

cathode plasma beam, and vapor deposition tech- niques. The files of the RMIC list 9 commercial and

5 research sources and more than 80 papers on

growth techniques since 1965. These papers do not include the large number of proprietary or in-house

reports of a number of commercial producers.

Papers A-1, B-2, B-7, B-8, B-11, B-14, E-1, and

E-14 of the International Conference on Crystal

Growth, 1967 [5] concern either sapphire or ruby,

and a great deal of information is available only in contract reports to the Air Force Cambridge Research Laboratory, Office of Naval Research, and

to other government agencies. The long-standing

program of research on "Crystal Growth and

Characterization" of the National Bureau of

Standards was changed on July 1, 1966, to a program

of research on "Crystal Growth and Characterization and High Temperature and Laser Materials."

W. Class et al. of the Materials Research Corporation have reported extensively [29] on electron- beam float-zone processing of aluminum oxide.

yttrium aluminum garnet, and multiply doped gar- nets. Large single crystals of sapphire have been

grown that are optically clear, have a typical

dislocation density from 10 to lO"* cm"^, and are completely clear of lineage boundaries. A typical

chemical analysis is 20 ppm Fe, Ni, 10 ppm Cr, Cu,

less than 10 ppm Ti, Si, and B, less than 5 ppm

Sn, and less than 1 ppm Mn, Mg, Pb. The electron- beam float-zone technique uses surface tension

to support a molten zone that traverses the crystal.

The purification arises for the most part not from

zone refining, but from selective distillation of

higher vapor pressure impurities. This is an advantage for sapphire, but a distinct disadvantage for ruby in which one obtains a radial distribution or coring of the chromium.

3.6. Hollow Cathode Plasma Growth of Laser

Host Materials

Yet another important technique in crystal

growth of oxidic materials is the hollow cathode

plasma beam process described by W. Class et al. [29] and by W. Class [30]. The advantages of this process are that one can operate at pressures up

to 3 Torr with virtually any gas and without any

thermionic electron emitter that might serve as a source of contamination. Float-zone purification

and single crystal growth of yttrium aluminum gar- net, yttrium aluminate, erbium aluminum garnet,

erbium aluminate, and the oxides of yttrium,

erbium, thulium, lanthanam, zirconium, titanium,

and aluminum have been carried out with this type

of hollow cathode. The crystalline perfection

appears quite good under crossed polaroids; however, this technique is just started, and there are no data on purity and perfection.

3.7. Czochralski Growth of Laser Host

Materials

Charvat et al. [31] and Keig et al. [32] of the

Crystal Products Division of the Union Carbide

Corporation, Indianapolis, Indiana, have reported

extensively on the Czochralski growth of ruby and

Nd^+-doped YAG. The most important characteristic for a laser host material is the optical

quality, and the purity of the starting material

is of utmost importance since impurities can precipitate during cooling to form minute but well

dispersed inclusions. These are difficult to detect,

but are a real problem since they cause a localized

absorption of optical energy under high-power

^-switched light operation and can lead to damage

within the crystal. The most informative techniques

for optical quality are Schlieren photography and

Twyman-Green interference. Good quality laser ruby has less than 1 fringe per cm of rod length.

202
where these fringes are due to strain or to uneven

distribution of solute. Crystals have been grown that are 5 cm diameter by 30 cm long, with no misorientations that can be detected by chemical

etching or the Schultz-Wei technique. The dis- location densities measured on the basal and prismatic planes are 10^ and 10^ cm"^, respectively.

The starting material for Czochralski growth of

ruby has 20 to 30 ppm of cation impurities. Ehmination of rare-earth impurities is of extreme importance for Nd^^ YAG, since these have optical

absorption tails that extend into the infrared region.

Faceting and core formation are also a problem

with this material.

4. Summary

The introductory lecture at the Santa Fe Conference on Radiation Effects in Semiconductors

(to be published by Plenum Press) was presented

by Dr. J. W. Corbett, General Electric Research

Laboratories, Schenectady, New York, on the

subject of "Important Unanswered Problems."

Dr. Corbett used three shdes as a part of his lecture

that read: What are the impurities; is that all the

impurities, and are you sure? It is felt that these

questions are equaDy applicable to the purity and

perfection of oxide materials; that solid state chemistry is as important as solid state physics; and that crystal growth and characterization are as important as research on oxide materials. Progress

in understanding the fundamental properties and

the ultimate range of properties depends, therefore, on a continued improvement in the purity and

perfection of research specimens.

5. References

[1] The author is indebted to Mr. T. F. Connolly, Director, and to Dr. H. S. Pomerance of the Research Materials Information Center for a great deal of assistance in assembling the information for this survey. Solid state crystal producers and users are urged to contact the RMIC,

Bldg. 3001, Oak Ridge National Laboratory, P.O. Box X, Oak Ridge, Tennessee 37830 for data sheets to record

their product or crystal desires.

[2] Engineering Properties of Selected Ceramic Materials,

1966, compiled and edited by BatteUe Memorial Institute (American Ceramic Society, 1966).

[3] The Growth of Oxide Single Crystals from the Flux Melt, by E. A. D. White: High Temperature Techniques, by

R. G. Bautista, in Techniques of Inorganic Chemistry,

vol. 4 (Interscience Publ., 1966).

[4] Report MAB 229, The Characterization of Materials, Materials Advisory Board (1967).

[5] Proc. Intern. Conf. on Crystal Growth, J. Phys. Chem.

Solids, Pergamon Press (1967).

[6] Z. Rabenau, Chemie Ingenieur Technik 36, 542 (1964).

[7] C. T. Butler and J. R. Russell, SoUd State Division, Oak

Ridge National Laboratory, are in charge of the program

of research on MgO.

[8] Y. Chen and W. A. Sibley, Phys. Rev. 154, No. 3, 842 (1967).

[9] R. W. Davidge, J. Matl. Sci. 2, 339 (1967). I am indebted

to Dr. Davidge for this figure.

[10] Y. S. Park and D. C. Reynolds, J. Appl. Phys. 38, No. 2, 756 (1967).

[11] Minnesota Mining and Manufacturing Company (3M),

Dielectric Materials and Systems Division, 2501 Hudson

Road, St. Paul, Minnesota 55119. I am indebted to Dr. Robert D. Carlton for the photograph and table of spec- trographic data.

[12] F. Galtier and R. Collongues, C. R. Acad. Sc. Paris C264,

87 (1967).

[13] R. J. Gambino, J. Appl. Phys. 36, 656 (1965).

[14] G. E. Pynchon and E. F. Sieckmann, Phys. Rev. 143, No. 2, 595 (1966).

[15] R. L. SprouU, W. C. Dash, and W. W. Tyler, Rev. Sci.

Instr. 22, 410 (1951).

[16] Personal communication with N. Quinn, W. & C. Spicer,

Ltd.

[17] F. K. McTaggart, Nature 199, 4981 (1963).

[18] R. Ruh and H. Garrett, J. Am. Ceram. Soc. 47, No. 12, 627 (1964).

[19] A. B. Chase and J. A. Osmer, Am. Mineralogist 51, 1808

(1966). [20] A. B. Chase and J. A. Osmer, J. Am. Ceram. Soc. 50, 327

(1967). [21] T. Niemyski and W. Piekarczyk, J. Crystal Growth 1, 177

(1967).

[22] J. R. Drabble and A. W. Palmer, J. Appl. Phys. 37, 1778

(1966).

[23] B. A. Smith and I. G. Austin, J. Crystal Growth 1, 79 (1967).

I am indebted to Dr. I. G. Austin for permission to use

the photograph.

[24] L. B. Robinson, W. B. White, and R. Roy, J. Mtris. Sci. 1, 336 (1966).

[25] A. T. Chapman and G. W. Clark, J. Am. Ceram. Soc. 48,

494 (1965). I am indebted to G. W. Clark for the drawing and photograph.

[26] A. T. Chapman, School of Ceramic Engr., Ga. Inst, of Tech.,

Atlanta, Ga. 30332, personal communication.

[27] C. B. Finch and G. W. Clark, J. Appl. Phys. 36, 2143 (1965); 37, 3910 (1966). I am indebted to G. W. Clark for this photograph.

[28] R. C. Linares, J. Phys. Chem. Solids 28, 1285 (1967).

[29] W. Class, H. R. Nasor, and G. T. Murray, Crystal Growth,

J. Phys. Chem. Solids (Pergamon Press) 1967; AFCRL67-0723 report (unpublished).

[30] W. Class, Research/Development (September 1967).

[31] F. R. Charvat, J. C. Smith, and D. H. Nestor, Crystal

Growth, J. Phys. Chem. Sohds (Pergamon Press) (1967).

[32] G. A. Keig, D. H. Nestor, and P. E. OttenV S'RCR-67-5, Annual Technical Report, Office of Naval Research

(unpublished).

203

The Growth of Oxide Single Crystals By Chemical Transport

Robert Kershaw and Aaron Wold

Department of Chemistry and Division of Engineering

Brown University, Providence, R.I. 02912

The growth of relatively large ferrite single crystals (5mm on an edge) was accomplished by an improved chemical transport procedure. Magnetite crystals were analyzed and found to have the composition Fe2.9904. Dry hydrogen chloride was used as the transporting agent and there was no ap- parent evidence of chloride contamination in the final crystals.

Key Words: Chemical transport, crysted growth, ferrite crystals.

Chemical transport reactions have been used [1]' to prepare single crystals of iron (II) diiron (III) tetroxide (magnetite) and other ferrites which comprise the majority of all known ferrimagnetic materials. Previously. Smiltens [2] had prepared well

characterized iron (II) diiron (III) tetroxide crystals

using a modified Bridgman-Stockbarger technique.

The crystals were grown under a CO2 atmosphere

and then CO was introduced and the ratio of these

gases changed gradually as the temperature was

varied during the coohng process. This maintained the proper oxygen pressure necessary for the preparation of stoichiometric crystals. The

techniques of flame fusion [3] and hydrothermal

growth [4] have also been used to grow ferrite crystals, but their quality, in general, has usually

been poor. In addition, a cobalt ferrite crystal was

grown by Ferretti et al. [5] from the melt at 1600 °C

under an oxygen pressure of 790 p.s.i. Chemical

analysis of a portion of the crystal gave a ferrous

ion content of 1.3 percent. Kunnmann et al. [6J reported that the simple ferrite spinels crystallize

readily from the ternary flux system MO-Na^OFe203. Crystals of CoFe204 were grown from a mixture of 1.5 moles of CoFe204 and 1 mole of

Na2Fe204 utihzing a platinum wire cold finger.

The crystals grown were approximately 1 inch in diameter and inch thick. However the crystals were not dense and appeared to have many defects. Nickel ferrite crystals were also grown by Kunn- mann et al. [7] from sodium ferrite flux by a modified

Czochralski method. The crystals were analyzed

and corresponded to Ni1.01Fe2O4.01. All of the latter methods were difficult to carry out and the results hard to reproduce.

' Figures in brackets indicate the literature references at the end of this paper.

Relatively pure crystals of magnetite and the

other ferrites can, therefore, best be grown by the

method of chemical transport. In this method [1, 8] the charge material reacts with the transporting

agent to form a more volatile compound. This

vapor diffuses along the tube to a region of lower

temperature, where some of the vapor undergoes

the reverse reaction. The starting compound is reformed and the transport agent is liberated. The

latter then diffuses to the hot end of the tube and

again reacts with the charge. Under the proper

conditions the compound may be deposited as

crystals. This may be illustrated with Fe304 where

the transport using HCI as the transport agent occurs by the reversible reaction

Fe304 + 8 HCI — FeCL + 2 FeCls + 4 H2O.

The above reaction is illustrated in figure 1, and

may also be used to prepare crystals of other

ferrites such as NiFe204, CoFe2 04.

Experimental Procedure

The pure starting compounds were prepared

by the method of Wickham et al. [9]. In this method

precursors are reacted to form the pyridinates

M3Fe6(CH3C02)i7030H • 12Py: and then careful

decomposition of the pyridinates results in well- crystallized, stoichiometric ferrites.

T= IOOO°C T = 750°C

CHARGE ZONE GROWTH ZONE

Phci = lOmm AT 300°K

Fe 304 + 8 HCI ^2FeCl3 + FeCl2 + 4H2 0 Figure 1. Transport process.

205
Figure 2. Vacuum system used to prepare transport tubes. The transport tube used for single crystal growth

is prepared by closing one end of a ten inch length

of 13 mm. I.D. silica tubing and sealing the other

end to an 18|9 vycor ball joint. After the tube has

been cleaned with aqua regia and dried a one gram

charge of ferrite is added through a long-stemmed

funnel, and the tube is connected to stopcock on

the vacuum system shown in figure 2. When the

pressure in the system has been reduced to less than 10"^ Torr the sample is outgassed by heating

to about 300 °C. Stopcock C is then closed, and HCl

gas is let in to a pressure of 10 Torr, as measured

with the mercury manometer. Stopcocks A and B

are then closed, the joint between them separated,

and the transport tube sealed off to a length of 8

inches with an oxy-hydrogen flame. This tube is then centered in the transport furnace and the

temperature controller themocouples are placed

at its ends. A Hevi-Duty MK2012 furnace, rewired

to give two independently controllable temperature

zones was used in these experiments. It was

found that a two cm. thick firebrick baffle was

needed at the junction of the two zones in order to obtain a satisfactory temperature profile in this relatively large diameter (6 cm.) furnace. The temperature of the empty, or growth zone is raised to 1000 °C, while the end of the tube with the powder,

or charge zone, is left at room temperature. This

back transport lasts for 24 hours, and cleans the

growth zone of stray nuclei by causing them to transport to the charge zone, where they cannot

interfere with crystal growth. The growth zone temperature is then lowered to 750 °C and the

charge zone raised to 1000 °C. Transport is allowed

to proceed for 10 days, after which time the charge zone temperature is reduced to 750 °C. In about one hour, when equilibrium has been re-established,

both heaters are turned off^. When the furnace has

cooled to room temperature (usually overnight)

the transport tube is removed and opened.

Properties

Crystals grown by chemical transport are well- formed octahedra, with shiny faces measuring up

to 5 mm on an edge. X-ray patterns of the groundup crystals were identical to that of the starting

material. The results of chemical analysis of a

typical magnetite run are given below. There was

no evidence of chloride in the final transported

crystals, although traces might be expected.

Total iron Iron (II) Formula

Found (Powder)

Found (Crystal)

72.36%

72.66% 72.83%

24.12% 23.78%

23.65%

Fe" Fe'" Oa r II p III Q

References

[1] Z. Hauptman, Czech. J. Phys., B12, 148 (1962).

[2] J. Smiltens, J. Chem. Phys.. 20, 990 (1952).

[3] R. Lappa, Przegland Telekomunikacyjng 31, 229 (1958).

[4] N. Yu Ikornikova, Dokl. Akad. Nauk. USSR, 130,610(1960).

[5] A. Ferretti, R. J. Arnott, E. Delaney and A. Wold. J. Appl.

Phys.. 32, No. 5, 905 (1961).

[6] W. Kunnmann, A. Wold and E. Banks, J. Appl. Phys., 33,

No. 3, 1364 (1962).

[7] W. Kunnmann, A. Ferretti and A. Wold, J. Appl. Phys., 34,

No. 4 (Part 2), 1264 (1963).

[8] R. Kershaw and A. Wold. Accepted by Inorganic Syntheses.

[9] D. G. Wickham, E. R. Whipple and E. G. Larson, J. Inorg. Nuel. Chem., 14, 217 (1960).

206
Panel Discussion

A. S. Nowick, Columbia University, Chairman

R. E. Carter, General Electric Company

J. H. Crawford, Jr., University of North Carolina

A. B. Lidiard, United Kingdom Atomic Energy Research Establishment, Harwell

D. H. Whitmore, Northwestern University Y. Haven, Wake Forest University

J. B. Wachtman, Jr., National Bureau of Standards

The purpose of the panel was to summarize and

comment upon the status of the field of mass trans- port in oxides. It was hoped that the discussion

would focus attention on important new ideas that have arisen in the conference, and would also point out possible avenues of approach that had been

missed or were not sufficiently emphasized earlier

in the conference. To accomphsh this objective,

the chairman selected a number of key questions. In each case, at least one member of the panel

presented the principal discussion, following which

the other panel members and the audience contributed to the general discussion. In what follows, we shall present the questions posed, the principal

discussions, and selected points from the general

discussion.

Importance of a Better Fundamental

Understanding

Question:

Is a better understanding of the fundamentals

of ionic mass transport in oxides reaUy needed in modern technology?

Principal Discussion (R. E. Carter):

The present question can be approached by considering three questions raised by Dr. R. Thomson

of ARPA at the start of the conference:

1. What is the state of our knowledge of ionic

transport in oxides?

2. What should be the goal of work on metal

oxides?

3. Who is going to do the required work?

What then is the state of knowledge of transport

in oxides and how useful is this knowledge to modern technology? Historically most of the work

in oxides has been done by what might be called

high temperature physical chemists, modern

ceramists and metallurgists. These people have

by and large gathered data to enter into analyses

which were quite approximate (simple mass action

equations, isolated defects, and activity coefficients of one). The motivation has been to understand

various processes such as metal oxidation, sintering, electroconductivity, plastic deformation, and

chemical synthesis. In general, to understand these

processes has not required data of high precision.

Also, as Professor Birchenall pointed out, many

of the measurements have been made on commercially available material because these are the

materials in which it is desired to understand the processes. These investigators have been fairly

successful in meeting the over-all goals that were

set up. It may then be said that our understanding

of technologically important processes such as

sintering is reasonably good. The principal problem

in sintering is one of models and, in the case of oxides on which the better transport data has been

gathered, understanding is not limited by the

absolute accuracy of the available diffusion coeffi- cients. Clearly this is not the case for aU ceramic

materials and an excellent case may be made for better transport data on oxides of specific interest. There seems to be little likelihood, however, that such data will improve our general understanding

of the important processes involving metal oxides.

Dr. Thomson's second question, the goal of work

on ionic transport in oxides, is related to the use to be made of the data. Several speakers at this con- ference have presented strong arguments asking for more accurate mass transport measurements both

in the intrinsic and extrinsic ranges. The attainment

of this accuracy will do much to improve the liaison between people interested in technological aspects

of diffusion and those interested in the more fundamental studies, which should be one of the goals of any diffusion study in oxides. However, many of the

processes that we wish to understand are not really

controlled by ionic motion, but are strongly influ- enced by other aspects, such as interfaces and grain

boundaries.

I think, therefore, that investigators working with

oxides now appear to be at a crossroads. They can

refine their techniques, work on better characterized

materials, and develop better theories of the defect

state. They can then hope or assume that this knowledge will be applicable to commercial materials and

processes. Alternatively they can examine their

207
processes directly and look at interfaces and grain

boundaries which may be controlling the rate of the

ionic transport of interest, an approach which is often ignored by those concerned primarily with

ionic motion. The greatest progress would appear to come from exploring both of these approaches,

which should ideally compliment one another.

Finally, who is going to do the work which wiU

permit a better understanding of oxides? The technologists, the ceramists and metallurgists, will no

doubt continue to work on materials that, by physicists' standards, are impure and poorly characterized, and they will probably move into more

complex areas like surfaces, interfaces, and grain

boundaries. The physicist, who makes precise

measurements and develops theories, will move

from the halides to the oxides. The rate of this move

wiU be approximately proportional to the availability

of high quality specimens. In this respect, it is clear

that "nothing succeeds hke success." Dr. Cleland

stated that he had more requests for MgO than any

other oxide. At the present time, MgO is probably

available in the highest purity and crystal perfection

of any of the oxides and probably better work has

been done on MgO than on any other oxide. While

the strength of the physicist is his ability and

interest in making precise measurements and de- veloping exact theories, his weakness is his lack of

interest in making good specimens. By and large he

waits until good specimens are provided. As ex- pressed by Werner Kanzig, "first come the crystals,

then the good measurements, and then the theory."

General Discussion:

Several people felt that Dr. Carter had overrated

the state of our knowledge and the usefulness of

presently available measurements for technological

purposes. The extremely poor reproducibility of

diffusion data was particularly cited. These are so poor that, with much of the data it becomes impossible to maximize diffusion-controlled properties.

Examples were also given of our inability to use

refractory metals for high temperature structural applications because of inability to protect ^them

from oxidation. Also, the processing of electronic

ceramic materials (e.g., ferrites, ferroelectrics)

is in a poor state because of inability to control

sintering and grain growth. Further emphasis was

also given in the discussion to the importance of

processes other than mass transport that must also be better understood, e.g., grain growth during

sintering, adhesion of scale, etc. Such processes

will undoubtedly remain the limiting factors in technological processes unless their understanding

is also advanced.

The question of proper characterization of the

specimens being studied was also discussed.

Obviously, only very pure or carefully doped specimens allow for a meaningful study. It is senseless

to try to develop a complex theory from precise

data if these data apply only to the particular specimen which was studied. A complete chemical

analysis is the minimum characterization which is now acceptable. With this analysis there should

be given the statistical distribution of the impurity

contents.

Value of Alkali Halide Work

Question:

What can we learn from all of the past work on

the alkali halides, and can we be misled by applying

these results too directly to the oxides?

Principal Discussion (J. H. Crawford, Jr.): As to the desirability of taking alkali halides as our point of departure in investigating oxides, I would like to make a few general remarks. Firstly,

there is a tremendous inheritance available to us from the long standing research program on alkali halides which dates back to the early thirties. This

inheritance is not only valuable to investigations

of defect behavior in oxides but has profoundly

influenced the development of solid state chemistry

and physics as it pertains to non-metalUc systems.

We cannot ignore it. The concepts of point defects,

composite imperfections, the mechanisms of their

creation thermally and by radiation, the role of

defects in diffusion processes, solid state reactions,

etc., we largely owe to studies of imperfections in alkali halides.

A second point concerns the importance of trace impurities in determining physical behavior of

non-metals: This has been vividly demonstrated

through investigations of alkali halides. Radiation

response (the yield of color centers), defect aggregation processes and related phenomena have been

shown to depend sensitively upon the presence of

certain types of impurities in very small concentration. Without the availability of high-purity alkali halide crystals which could then be doped to the

desired level with the desired impurity, it is doubtful

that this impurity sensitivity could have been ex- plored. At this point Hve should acknowledge that

in many systems impurity sensitivity may not be

nearly so great as in the alkali halides and that some equally dominating structural characteristic

may be the focus of deviations from intrinsic behavior. For example nonstoichiometry which is notorious in the oxides of transition metals may be

the main point of interest. Nevertheless, the importance of alkali halides as a simple model for defect behavior and mass trans- port in binary systems is clearly evident, whatever

ultimate modifications to such a model might be

necessary to make it applicable to a given oxide

208
system. As a matter of fact, modifications to the

simple model are always necessary in order to

utilize the body of knowledge available. To use a

specific case, let's consider how the concepts are transferred from alkali halides to MgO. Here we

must examine differences as well as similarities. The predominant distinction is that MgO is a di- valent binary system with the consequence that the

binding energy and the energies of formation and

motion of cation and anion are considerably higher

than for an alkaH hahde with the same lattice spacing. This means among other things that for a given impurity content one must go to a much higher

temperature range than in an alkali halide to reach

the domain of intrinsic (as contrasted with impuritysensitive) behavior as applied to mass transport.

Another consequence has to do with possible

stable impurity-defect complexes: Whereas a divalent cation in an alkah halide introduces a

single vacancy which can be bound to an impurity

to form a dipolar complex, a cation impurity with one surplus charge in MgO can account for only

one-half of a cation vacancy and the number of

stable and metastable complexes is large, e.g., an impurity-vacancy complex still bearing a negative

charge, and neutral I2V complexes with a variety

of configurations. It is interesting to note that the I—V complex by virtue of its charge can drift in an

applied electric field and contribute to the ionic

conductivity. Another difference is radiation re- sponse: Sibley and coworkers at ORNL have shown

that although color centers can be readily created

in MgO by fast neutrons and 1.5 Mev electron ir- radiation, their creation is evidently caused by a knock-on process rather than a radiochemical one

since x rays and y rays are incapable of producing

F-centers. On the basis of the difference in valence

alone it is evident that we face a much greater range

of complexity in both number and response of

imperfections. This should not necessarily be taken

as a disadvantage since by proper use of some ele- ments of this complexity it is possible to gain a deeper insight into the nature of the defects and

their interactions. Nevertheless, it must be admitted

that getting a specimen of MgO sufficiently pure to lower the intrinsic range of behavior into an experimentally convenient temperature range may be a

difficult task. General Discussion:

It was pointed out that in the study of nonstoichiometric oxides one might encounter new phe- nomena not previously seen in the alkali halides.

One such effect is the ordering of defects to pro- duce distinct phases of material, as described in the paper by Libowitz. Another telling point made

was that alkali haUdes are essentially completely

ionic whereas many important oxides and oxide

systems are virtually covalent. Therefore, it would

be very advantageous to have another well understood prototype material at the other end of the

scale. Unfortunately with the best of these, i.e., an elemental material such as silicon or diamond,

it is difficult to make a connection with the proc- esses expected in a binary oxide. Nevertheless a predominantly covalent III-V or II-VI compound

might be helpful to use as a guide.

Importance of Theory

Question:

What role can theory play in the development of

the subject of mass transport?

Principal Discussion (A. B. Lidiard):

One must consider two principal parts of the

theory (i) the underlying mechanics of the defect

solid, its energy levels and states, the configurations of individual defects and so on and (ii) the

statistical mechanics of defect solids, generally in or near thermodynamic equilibrium. Since models

provided by this second part of the theory have

formed the basis for most of the discussion at this conference I will deal with this part first. Statistical mechanics of systems in thermodynamic equilibrium. Here one is interested in partial

thermodynamic quantities, dependence of composition upon oxygen partial pressure, etc. Past

discussions [1, 2] ' have generally been in terms of

model systems of non-interacting point defects —

the direction in which the subject was very successfully launched many years ago by Wagner. The

extension of these models requires us first to specify

the interactions between defects and then to cal- culate their effect upon the statistical properties.

At long-range the dominant interactions between

defects bearing a net charge are undoubtedly Coulombic. The only way of handling these which is sufficiently simple to be easily and universally

applied is that provided by Debye-Huckel theory: more elaborate and accurate calculations for simple

systems show that it is only qualitatively cor- rect [3]. However, since, as Prof. Slifkin's review

(this conference) showed, these Coulombic defect

interactions are qualitatively important at concentrations even as low as 0.1 atomic percent they

must certainly be important at 5 or 10 percent (as

e.g., in nonstoichiometric FeO, UO2 etc.) and it is certainly desirable therefore to make much wider

application of the Debye-Hlickel description even

though it may be only approximate.

The evaluation of defect interactions at small

separations is more difficult and depends strongly

upon defect structure and details of the inter-

' Figures in brackets indicate the literature references at the end of this discussion.

209
atomic interactions. Evidently when the defects

have a complex or unsymmetrical structure (as in the Roth model for Fe vacancies in Fei-j-O) this interaction depends on the mutual orientation

of the defects. However, even for the simple ex- ample of two vacancies in a metal the interaction

depends strongly on the direction of the axis of the

pair relative to the crystal axes; it is not a simple

function of separation as suggested by continuum

theories [4]. It is unlikely that theory can provide

much insight into the interactions occurring at small

separations in the absence of reasonable models for the substances of interest; these are only likely

to emerge from a general and basic study of their

physical properties. However one aspect of the

interaction of complex and unsymmetrical defects can be handled without difficulty and that is "hindrance" effects i.e., the way in which one defect

by its extension prevents another defect from occu- pying a certain number of neighbouring sites. The

influence of these hindrances on the predicted

composition isotherms was illustrated by the work

of Libowitz on Fei-^O. I would interpret this work

(which neglects other interactions) as showing the

importance of hindrances but not necessarily the

unimportance of those features omitted from the

calculation (e.g., interactions at short and longrange). It is implied by these remarks that it is adequate to regard the oxide as a perfect lattice containing a statistical distribution of atomic

defects — a distribution which can be determined by

application of the methods of statistical mechanics.

Of course, the model of interacting defects allows

the possibility of cooperative ordering phenomena —

as was recognised roughly twenty years ago, particularly by Anderson. It has been used by others

since, notably by Libowitz for metallic hydrides [5]. The Anderson treatment was based on the simple

Bragg-Williams approximation and the theory of

physically simplified order-disorder systems (e.g., the Ising model) has now become mathematically

very sophisticated. It is possibly time that theorists

interested in the field of cooperative phenomena

turned their attention to more elaborate models.

The occurrence of MagneH phases seems to pre- sent a rather striking problem for the theory of

these oxides and one where it is not, however, evi- dent that the atomic defect model is a good starting

point.

Statistical mechanics of non-equilibrium systems;

diffusion etc. Considering kinetic processes, one

sees first that for dilute systems of defects one has

a fairly good theoretical framework already avail- able. This was reviewed by Manning (this conference). The scheme is sufficiently general to be useful but it needs to be filled out experimentally by

the determination of parameters such as the various

defect jump rates and activation energies by which

the model is specified. This is being done successfully in some metals and alkali halides and in principle could be done also for the closely stoi- chiometric oxides like MgO etc. When one turns to concentrated systems (high

defect concentrations, chemical diffusion couples,

etc.) there is very little to go on. Manning has made

many of the extensions of the kinetic theory nec- essary to deal with chemical diffusion in metals.

This gives hope that a similar development of practical value could be made to oxide systems.

In this situation perhaps the most that can be

done is to relate experimental results to a general

phenomenological scheme adequate to distinguish

clearly the different diffusion coefficients which

arise (as e.g., in oxidation, interdiflfusion of two

oxides, tracer diffusion etc.). Such a scheme is provided by the phenomenological equations of

irreversible thermodynamics [2]. However, it should

be noted that this scheme of equations can make

only very limited physical predictions in the ab- sence of either experimental or theoretical knowledge of the behaviour of the phenomenological

coefficients as functions of concentration. Howard

and 1 have elsewhere [6] discussed the form of these

phenomenological coefficients in dilute systems in both ionic compounds and metals. Although for more

concentrated systems the scheme cannot predict

much that is physically new, it may be very useful for keeping clear the differences in the nature of the

diffusion coefficients measured in different experi- ments. While this is true for binary systems it is even more true for ternary systems on which there

appears often to be an interesting limitation (Cooper,

this conference). One should recognise however

that the scheme of phenomenological equations one

is concerned with here introduces six independent

phenomenological coefficients for an isothermal

ternary systems, each of which is a function of composition and temperature, and that it is very

unHkely that these can all be determined by ex- periment—simply because there are not enough

independent experiments which one can do. Further progress depends upon being able to obtain

insight into the form and magnitude of these

coefficients from defect models.

States, Energy Levels, etc. of Individual Defects. Although this subject logically preceeds that of

the statistics of systems containing defects I deal

with it last since it was discussed less frequently

at the present conference. In this part of the subject we are almost always concerned with very low

concentrations of defects when we are using spectroscopic methods; these methods (optical — es- pecially on zero-phonon and other sharp lines,

e.s.r. etc.) give important information about the

210
structural and electronic symmetry of the defects being studied (cf. the papers of Henderson and

Kemp, this conference). The theory of the electronic

states of these electron-excess colour centres in the alkahne earth oxides (F-centers and their

aggregates especially) has been taken over from

the alkaU halides with little modification of principle and is very successful. Much less has been

done on the electron-deficient centres (trapped

holes) and rather little on the states of defects in transition or actinide oxides except for the study

of impurity states. The complex structures inferred

for the iron vacancy in Fei-jO and the oxygen in- terstitial in UO>^x from the long wavelength

neutron scattering studies of these materials

have not received a fundamental explanation.

Turning now to the calculation of defect formation and migration energies I would say this problem does not look very straightforward. Boswarva's

results presented to this conference showed a consistency of different variants of the ionic model

for the alkali halides that was lacking for the

alkaline earth oxides. This is perhaps partly en- hanced by the unreliabiUty of the necessary but

quite basic experimental data (e.g., compressibilities). It is likely that the models used for calculating

defect migration energies, etc. will only give

useful results if they are developed as part of the

general sohd state physics of these materials. Here

the work on lattice vibrations which Slater reviewed

can play a very important part. One can see in particular that good shell models have been built up

for many materials from neutron scattering experiments but have not yet been much used for defect

calculations although the way to do so is pretty

clear. It would help considerably to know the pre- dictions of these models for the alkahne earth

oxides, UO2 etc. In oxides like Fei-^O and UO2+X it is much more

important to know the electronic and charge states of the defects. Thus in considering the interstitial configuration in UO2+X it is important to know

whether the electron holes are locahsed near the

oxygen interstitial. From measurements of partial

thermodynamic quantities and from electrical measurements on nonstoichiometric UO2 one can

say that it is certain not both holes are so locahsed.

It is not, however, clear that both of them are nonlocalised (inferences have generally been made by

comparison with models which omit interactions).

If one of the two holes is locahsed on or near the

oxygen interstitial, then a much more elaborate

calculation is needed: the atoms whose interaction one is attempting to deal with, are thus not in the same charge states as those which determine the

lattice vibrations, elastic constants etc. Familiar

examples of the importance of electronic state in

determining the equilibrium atomic configuration

of the center are provided by the alkali halide hole centers Vk-, H etc., which have (110) axes of symmetry. Fundamental discussions of the depend- ence of configuration upon the electronic state of

simple vacancy centers in a fuUy covalent system

(diamond) have been given recently [7].

General Discussion:

The importance of three-body forces (cf. the

departure from the central-force Cauchy relations

for MgO) was mentioned as a possible source of

difficulty in defect calculations. Lidiard's reply re- ferred to Slater's view, presented to the conference,

that noncentral forces are not very important for the overall dynamical properties, and suggested the same may be true of defect properties. Explicit

attention to the role of three-body forces in defect

properties has thus far only been given for the

inert-gas soUds [8]

.

References

[1] For an introduction see P. G. Shewmon, Diffusion in Solids, Chap. 5, McGraw Hill Book Co., New York (1963).

[2] For a fuU and modern account of all aspects of mass trans- port in solids see Y. Adda and J. Philibert, La Diffusion dans les Solides. Presses Univer staires de France. Paris

11966). For oxides see especially II, Chap. XIX.

[3] Review by A. R. AUnatt, Adv. in Chemical Physics, 11,

1 (1967).

[4] J. R. Hardy and R. BuUough, PhO. Mag. 15, 237 (1967).

[5] G. G. Libowitz. J. Appl. Phys. Suppl. 33, 399 (1962).

[6] R. E. Howard and A. B. Lidiard, Reports on Progress in Physics, 27, 161 (1964).

[7] J. Friedel, G. Leman and M. Lannoo. Phys. Rev. 164, 1056

(1967) also M. Lannoo and A. M. Stoneham. J. Phys. Chem.

Sol., in press.

[8] See e.g. the review by G. L. Pollack. Rev. Mod. Phys. 36,

748 (1964): also R. Bullough, H. R. Clyde and J. A. Venables,

Phys. Rev. Letts. 17,249(1966).

Analysis of Diffusion Data

Question:

Why are data on mass transport in oxides generally so poor and what can be done to overcome

the obstacles to better measurements? Also,

what are the relationships between the various

types of diffusion data which are being obtained?

Principal Discussion I (D. H. Whitmore):

This conference has certainly pointed up emphatically the need to differentiate clearly between

the various kinds of diffusion coefficients being

evaluated for oxide systems. To date, many of the

diffusion experiments conducted on oxides have

been concerned with either: (a) the measurement

of the dififusivity of the cationic species in a "pure"

simple oxide, utilizing an appropriate radioactive

tracer of that species and a sectioning method or

a decrease-in-surface radioactivity method: or

(b) the measurement of the diffusivity of oxygen,

using an oxygen isotope exchange method which

211
may involve either sectioning the specimen after the diffusion anneal or sampling to determine the

composition of the gas phase surrounding the dif- fusion specimen as a function of time. It is mainly,

these experimentally-derived diffusion coefficients

which are of interest to the theorist because they

serve as a basis for comparison with the corresponding diffusivities which he derives from his atomistic

theory of diffusion processes in oxides.

On the other hand, the technologist concerned

with sintering, oxidation processes, etc., may be more interested in the chemical or interdiffusion

coefficient which is derived by analyzing the concentration profile in an oxide solid solution formedby interdiffusion between the two different component oxides forming the diffusion couples. This

latter diffusion coefficient is a phenomenological one and, as Dr. Lidiard has already remarked, is usually not related in a simple manner to the dif- fusion coefficients for the individual species in the

oxide solution being studied.

If we consider first cation tracer and oxygen ion

diffusion measurements that have been carried

out on simple oxide systems by different investi- gators, the reproducibility in the measured dif- fusivities is generally poor. In fact, past experience

indicates that the chance of one investigator's value

agreeing with another's (measured under pre- sumably identical experimental conditions) within

50 percent is fairly small. Since the reproducibility

of a given tracer diffusion coefficient in a metal

or alkali halide to within 2-5 percent is now a com- mon occurrence, adapting the experimental techniques utilized by investigators in these fields to the study of cation and oxygen diffusivities in oxides

might yield significant improvement in the oxide

data. It has been reported that Argonne investigators working with simple oxides have succeeded

already in achieving such reproducibilities for their

cation tracer-diffusion coefficients when such ex- perimental methods were employed.

The first serious problem facing the experimentalist undertaking the determination of the

cation tracer diffusion coefficient in a given oxide

is that of obtaining well-characterized oxide single

crystals of extreme purity for his measurements.

In contrast with the impurity levels obtainable in semiconductors and alkali halides, most oxides

available today must be judged to be relatively

impure materials. This may be a very serious

problem if the oxide of interest falls into Professor

Birchenall's "impurity sensitive" category, or less serious if it falls into his "impurity tolerant" cate- gory. Generally speaking, a precise knowledge of the

dislocation density, impurity content and distribution, mosaic structure, etc. for the oxide crystal

being studied is essential (although not usually

done) if a meaningful understanding of mass trans- port behavior is desired from the experiment.

Next, in conducting a diffusion anneal of the simple oxide specimen (a binary thermodynamic

system), the temperature, hydrostatic pressure and

one other thermodynamic variable must be fixed

in order to define unambiguously the thermodynamic state of the system (the oxide diffusion

specimen) being investigated. In this case, it is often convenient to choose the thermodynamic

activity of oxygen as the third independent variable

since it may be controlled experimentally by reg- ulating the composition and pressure of the gas atmosphere surrounding the specimen during its diffusion anneal. [To extend this consideration to the more general case of a multicomponent oxide

system containing r components, we note that this case requires that (r— 1) independent variables

be specified in addition to the temperature and

pressure. The problem of which variables should

be fixed in a chemical diffusion experiment involving

a multicomponent oxide is frequently not a trivial one and warrants more careful consideration than

has been given to it in several past diffusion in- vestigations on oxides.]

The annealing operation for the specimen is certainly one of the most vital steps in the diffu- sion experiment. Component parts of the heating

system like the furnace muffle, specimen container,

radiation shields, etc., as well as the surrounding

gas atmosphere, may serve to contaminate the spec- imen itself during the diffusion anneal. The precise measurement of the temperature of the anneal

and its duration is important, and the former

measurement may be a severe problem for the

experimentalist when the anneal is being carried out at ultra-high temperature. It is likely that, for those oxides classified by Professor Birchenall

as "impurity sensitive" or "borderhne, " the diffu- sion annealing ranges should have been extended

to appreciably higher temperatures in order to insure that intrinsic diffusivities were indeed

observed.

Measurement of the concentration profile of the

diffusing species in the oxide specimen following

the diffusion anneal has not received the attention

it deserves in many laboratories involved in oxide

diffusion experiments. Experience gained from

diffusion investigations in metals or alkali halides

has served to point up the value of determining

the oxygen isotope or cation tracer concentration

profile by either microtoming or precision grinding

layer by layer off the diffusion specimen. The

microtome sectioning technique might be used for oxides if a diamond knife were employed and the

precision grinding methods, common for sectioning metal diffusion specimens, can readily be

212
modified for oxide specimens. Short-circuiting dif- fusion paths (dislocation pipes, cracks, etc.) may

have a profound effect on the observed diffusion

coefficient and are best detected by autoradiographic examination of oblique sections through the

specimen.

Although Dr. Haven's discussion which follows

deals with the subject of interdiffusion in detail,

the measurement of a chemical diffusion coefficient

has several unique problems not dealt with in the

above remarks. One of these is the bonding to- gether of the two halves of the diffusion couple so

that there is no barrier to interdiffusion at this interface. This may be accomplished, for instance,

by hot pressing the two pieces together in an induction-heated graphite die. The determination of the

concentration profile may be undertaken with the

aid of electron microprobe scanning parallel to the

diffusion direction; this same technique should

also prove useful in revealing short-circuiting dif- fusion paths in the specimen.

In summary, then, the problems currently en- countered in diffusion experiments with oxides are

invariably more severe than those encountered

when deeding with metals or alkali haHdes. Some

of the factors which are undoubtedly responsible

for this present situation are: (1) the relatively

impure and poorly characterized oxide single crystals available: (2) the contamination and temperature measurement problems arising because the

diffusion anneals for the oxide specimens, of

necessity, have to be conducted at extremely ele- vated temperatures; (3) the imprecision, achieved

by the majority of the experimentalists who have

investigated mass transport in oxides, in the evaluation of the concentration profile.

Principal Discussion II (Y. Haven):

The use of different diffusion coefficients calls for some caution, because they can be defined

for different sets of gradients. A general discussion

of such relationships, based on irreversible thermodynamics, is given by Howard and Lidiard (Reports on Progress in Physics 27, 161 (1964)).

For a particular example one may decide to use one of the descriptions:

Ja = Lnid/jLAldx) + Lnid/JLBldx) + Lisid/jbcldx) (1) = L'uidcAldx) + L\2(dcBldx) + L[^(dccldx) (2)

jA={L;-L[.,){dcAldx) (3) = {Lu-L[,)idcAldx) (4)

where Ja is a current of A, the L's are coefficients

and dfjildx and dc/dx are gradients of chemical

potential and concentration, respectively. Equations (3) and (4) follow from eq (2) with the as- sumptions

dcAldx — — dceldx , dcci dx = 0

and dcAldx = — dccIdx , dcs/dx = 0

,

respectively.

Suppose A and B are isotopes and C is a second

component.

When the tracer diffusion is considered one finds

from (3)

D%=L[,-L[^=f{nlN)Da,i,,i (5)

where Ln and LI2 are of the same order of magnitude. The result contains a factor (nlN) for the

vacancy (or interstitial) concentration and a factor

/ for the correlations, /^defect is a diffusion coeffi- cient defined for the random motion of defects.

It should be remembered that irreversible thermodynamics itself does not give factors like / or

(n/AO, which can only be incorporated by specific

kinetic arguments. Equations like the Darken

eq (6) cannot account for these. One obtains rather simple relationships with the

chemical diffusion coefficients, in cases where

A and C have similar diffusion mechanisms, for example in an aUoy AC. A diffusion coefficient

described by eq (4) has the same structure as the

tracer coefficient (3). In such a case usually the

Darken equation gives a satisfactory description:

D = {DfNa + DSNi ){l + d\nyldln Ni) (6)

where the A^'s represent mole fractions. In the case of oxides, however, if A and C rep- resent cation and anion respectively, this proce- dure can be completely unsatisfactory, because the

two components may travel on different sub-lattices, or may be coupled by electric forces. For

example, in the interdiffusion of oxygen and cation a correlation factor is usually absent, also the

factor (n/N) may appear in a different context. In such cases one needs to go to the formula as developed by Schottky and Wagner (see paper by

Heyne) and formulae for interdiffusion often take

a form

D = 2DaDbI{Da^Db) (7)

which equation can be encountered in ambipolar

diffusion, or in the oxidation of metals.

If one component is slow, one may find for example

D = 2Da. (8)

213
Note the factor 2. When a single phase oxide is subjected to a change of oxygen pressure, the attainment of

equilibrium is governed by a diffusion coefficient

(4) which takes the form

D= (L;i-L;3)=£'defect (9)

Comparison of eq (9) with the tracer diffusion

coefficient of eq (5) enables one to obtain separately the concentrations and mobilities of defects. The latter information is to a certain extent lost, however, if the results are discussed in the lan- guage with chemical potentials as in eq (1), etc. (which may be the only feasible system).

For example, the rate constant for the oxidation

of metals is in Wagner's theory presented in the

form

1 300 1 f'^A<f) ^"'^~W7\ l^r \ (7 • riion neidfiA (10)

\ZA\ t Ne J 1x^(0)

where cr is a conductivity, njon and n^i are trans- ference numbers for ionic and electronic transport,

^a{o) and /i.^(^) are the electrochemical potentials

at the boundaries of the oxide phase, and the other

symbols have their usual meaning.

Only in the case that ngi is close to unity does the

integrand represent the ionic conductivity and can

it be replaced by a diffusion coefficient (hereafter

denoted by the symbol D, and representing the sum

of the diffusion coefficients of the ionic species

involved, if properly normalized).

In more elementary approximations of eq (10) the conductivities (diffusion constants) may be

independent of the position in the layer or rather

independent of fXA so that eq (10) can be integrated

in a straightforward way to

. ... A„,-£>(M.4(^)-At,(0)) (11)

leading to Darken-type equations (except for a

different type of summation of the D's). Such type

of equations may be useful for the discussion of

transport of oxygen through layers with a rather wide

range of composition, because in these cases

the diffusion coefficient may not be too strongly

dependent on compositions. If this condition

holds an equation for the chemical diffusion

coefficient D in the form

Z) =(!//)«£>* (a In a.4/d In C), (12)

where is the activity of component A, may be

quite appropriate, if a is a constant still to be de- termined for each particular case. However, if the D is dependent on composition an average D

is measured which may differ considerably from

D's measured in other ways.

Basically, eq (10) is derived with the assumption that the currents are throughout the layer the

same, which condition can be maintained because

the gradient of ixa can adjust itself to this condition. However, in order to use such experiments

for the determination of Z)'s, one must ascertain

that no barriers at the interfaces occur. Especially in such type of experiments large gradients

may occur at one boundary, which may produce

erratic results, because these gradients can be

very sensitive to disturbances.

Many so-called irreproducible results are due

to a lack of specific data on the thermodynamic

circumstances.

A Materials Program?

Question:

Would it be desirable to select two or more particular oxides to serve as prototype materials, and

to make interlaboratory comparisons on them? If so, which materials should be chosen, and how

pure and well characterized would they have to be?

Principal Discussion (J. B. Wachtman, Jr.): The uncertainty and lack of reproducibility which

are frequently mentioned in connection with dif- fusion data make it very desirable to have some means of checking on both the experimental tech- nique and on the specimens as possible sources

of variability. Interlaboratory comparisons are a useful technique for assessing the accuracy of a given measurement and are sometimes also useful

in helping to find the source of errors. For such

comparisons, measurements must be made on

the same specimeh or on specimens having the same properties within the accuracy desired. For

certain diffusion techniques, such as oxygen self- diffusion by gas sampling, it may be possible to use the same specimen in different laboratories,

but diffusion experiments usually require destruction of the specimen so that reproducible speci- mens are needed.

A reasonable degree of reproducibility has been

achieved in diffusion experiments on certain metals,

alkali halides, and silver halides and one might

consider the use of some of these materials as inter- laboratory comparison specimens. The temperature

range and the chemical environment (generally

oxidizing) required for studies of diffusion in oxides

do not favor the use of metals or halides, however.

It appears that oxide specimens must be used for

this purpose, but variation of composition and microstructure between nominally identical oxide speci- mens is thought to be one of the major sources

of variability in diffusion data on oxides. If a re- 214
producible oxide can be obtained, it might be used

to test the adequacy of experimental techniques

which, when proven sound, could then be used to search for the sources of variability in other oxides.

This proposed program would actually require a sequence of different oxide intercomparison speci- mens to correspond to the wide range of experimental conditions used for cation and anion diffusion

in the various oxides being studied. The whole

program would depend on producing one or more

oxides which are sufficiently reproducible to serve

as starting points.

The class of impurity-tolerant oxides discussed

by BirchenaU in this conference appears to offer the best prospects for such starting points. Dif- fusion in these oxides is dependent upon a very

high concentration of point defects which is controlled by fixing the departure from stoichiometry

through control of the activity of the major components or by fixing the amount of one additive at such a high level as to dominate other impurity

effects. On this basis, the following list of oxides

is suggested:

Material

85 Zr02+15 CaO

Fe,_xO

98.5 Th02+1.5 La.Os

MgO

Type and application

Impurity-tolerant; oxygen

self-diffusion Impurity-tolerant; cation

self-diffusion Impurity-tolerant; variable

mixed ionic and elec- tronic conduction

Impurity sensitive; self-dif- fusion of both components, foreign cation

diffusion.

The suggestion of calcia-stabilized zirconia as an

interlaboratory comparison standard for oxygen self- diffusion was based on the comparatively good

agreement in diffusion constants obtained by several different methods (see BirchenaU's paper). It

is known, however, that at 1000 °C the transport

properties change slowly with time as ordered microdomains form. This might seriously limit the use of

calcia-stabilized zirconia as a comparison standard

at temperatures below 1400 °C. Wiistite was sug- gested for comparisons of cation self-diffusion, be- cause it appears possible to fix a large departure

from stoichiometry by using a practically achievable

level of oxygen pressure, and because considerable

work has been done on this system with results which show good consistency for oxides. Thoria

containing lanthana or yttria was suggested as a

mixed ionic and electronic conductor whose per- centage of ionic conduction can be varied by varying

the partial pressure of oxygen; this behavior is discussed in Heyne's paper in the present confer- ence. Such a variable mixed conductor would be

valuable in assessing experimental techniques for measuring ionic conductivity for comparison with

diffusion data through the Nernst-Einstein relation. Heyne has discussed various experimental pro- cedures all of which place strict requirements on

electrodes as to their being reversible, or blocking

to ionic current, or blocking to the electronic current. A controllable variable mixed conductor would

permit checking some of the electrodes now used.

Finally, the use of MgO as an impurity-sensitive

comparison standard was tentatively proposed be- cause much effort is being devoted to making better

crystals (see the paper by Cleland) and a level of

purity and perfection of MgO may be achieved

which would permit its use to check experiments for contamination from supports and furnace atmos- phere which may be a serious problem at the high

temperatures used in diffusion studies in refrac- tory oxides. In particular, oxygen self-diffusion usually involves a surface layer only a few microns

thick and it is just this surface layer which is rather

difficult to protect from contamination in the range

of a few parts per million. It is frequently not ap- preciated how difficult it is to heat an oxide to 1800 °C without trace contamination of the surface. General Discussion:

This was a lively discussion in which many participants, while generally agreeing that inter- laboratory comparison specimens are desirable,

pointed out a number of difficulties associated with

the ones listed by Wachtman and suggested some

other possibilities:

(1) There is some evidence that surface ex- change of oxygen is the slow step for thoria

and stabilized zirconia. This affects the diffusion

profile obtained in a gas-soHd experiment and

would considerably complicate the analysis.

(2) The existence of time-dependent changes

in the zirconia-calcia system associated with

ordering of the cation sub-lattice was emphasized and the possibility was suggested that

this may also occur in lanthana-doped thoria

although no evidence was offered. The electrical conductivity depends on the grain size in calcia- stabilized zirconia which suggests a possible

dependence of oxygen diffusion rate on grain

size also.

(3) It was pointed out that a eutectoid reaction

and precipitation of magnetite occurs in Fei-^O

so that heating and cooHng the specimen would

probably cause changes. In reply, it was argued

that proper control of the activities of the com215
ponents of the system being studied, for sufficient time to achieve equilibrium, is an essential part

of a diffusion experiment in a compound. The

sensitivity of Fei-xO to change of temperature

and activities of components is thus a desirable

feature of a standard specimen for checking

experimental technique. This statement was

accepted, but the original objection was clarified by pointing out that irreversible structural

damage may occur in the process of precipitation

and reheating. Thus, even if the specimen was

equilibrated with respect to composition at the same temperature, structural damage might

remain and affect diffusion values.

(4) Cobaltous oxide was suggested as a better

possibility than ferrous oxide, although precipitation of a higher oxide is thought to occur

in Coi-j-0 as well as in Fci-j^O. This precipitation may not be an insuperable difficulty for either of the two oxides. The combination of choosing a small value for x and using a rapid

heating and cooling schedule may sufficiently

reduce structural damage.

(5) It was suggested that heavy doping of some

oxide other than zirconia or thoria be tried to produce an impurity-tolerant standard: magnesium oxide doped with 5 percent alumina was

suggested as an example. In reply, it was pointed

out that while soHd solution of AI2O3 in MgO

is thought to exceed 5 percent at 2000 °C, precipitation of spinel would occur below 1700 °C. The general idea of heavy doping to produce a

solid solution which is impurity-tolerant is attractive, but difficult to carry out because the

possible existence of solid solution below 2 percent has not been investigated for most oxide

systems. In those systems which appear to exhibit a large degree of solid solution the solid solution limit usually decreases rapidly with

decreasing temperature, causing precipitation. In other cases (such as calcia-stabilized zirconia),

the solid solution is metastable at intermediate

temperatures and formation of microdomains

occurs if time is allowed.

(6) The contamination of specimens during

heating can be reduced by enclosing the specimen

in a container of the same material. This may

be a useful procedure in some diffusion experiments, but does introduce an additional complication in oxygen exchange experiments.

(7) Calcium oxide was suggested as an alter- native to magnesium oxide because a material

having cations and anions of approximately the same size and polarizability is desirable for comparison with alkali halides. The best understood of the alkah hahdes are those having ions

of similar size such as NaCl or KCl.

Summary

The problem of mass transport in oxides may be

divided rather naturally into two parts. The first part deals with the fundamental lattice defects

present in pure (or, at least, well defined) single

crystals and their role in mass transport. Here, the

people studying oxides invariably look with admiration upon the experimental and theoretical achievements in similar studies of the alkali halides. There

is little doubt that the alkali halides provide a natural

starting point for the study of oxides because of the

comparative simplicity of those materials as well

as because of the accumulated information about

defects in them. On the other hand, investigators

of the oxides must be on the lookout for new phenomena, particularly those related to the higher

valencies as well as the more covalent bonding of

the oxides. As examples, the much greater deviation from stoichiometry in the oxides and the occur- rence of microdomains should particularly be cited.

If the lessons of the effects of purity in the alkali halides can be carried over to the oxides, it is clear

that one cannot work too hard in the direction of

improving the quality and purity of available oxide

materials. In fact, it is doubtful if the intrinsic range

of behavior can be observed for many of presently

available oxide crystals.

For an advance in theoretical understanding,

there is a great need to obtain good values for the

formation and migration energies of the various

predominant defects, in addition to information on

defect symmetries and electronic charge states. At present, there is no expectation that theory alone can give meaningful values for these quantities.

Rather, the experimentahst must supply the numbers which will then provide the basis for testing of theoretical models, the types of potentials to be

employed, etc. This seems to be the best approach

for advancing our theoretical understanding of

defects in oxides. Such a program then requires the

carrying out of such basic measurements as tracer

diffusion of both cation and oxygen, ionic conductivity, transport number measurements, and where

applicable, electron spin resonance and optical

absorption. In all cases, extreme care is needed

to fix the thermodynamic variables and to prevent

contamination when measurements or heat treat- ments are carried out at very high temperatures.

But all this care is not meaningful without good

starting materials!

The need for better materials has been mentioned

in the discussion at this conference, over and over

again. It was particularly interesting to note how

much debate arose over the proposal by Wachtman of four materials to serve as interlaboratory

standards. The objections raised, particularly emphasized the complexity of structural and phase

216
changes that may be expected when sohd solutions or highly nonstoichiometric oxides are chosen.

Another factor that might be added here is the

desire of solid-state scientists to make basic meas- urements on the purest and simplest materials, so

that good data will have "handbook" value, especially to theoreticians. In this connection, the pro- gram of the Oak Ridge National Laboratory to prepare MgO crystals of high purity is probably

most relevant. The availability of any one oxide of

high purity and perfection will inevitably stimulate more reliable data and substantial interlaboratory

comparison. In addition to the role of the national

laboratories in producing such materials, however,

industrial firms which produce or have a strong

interest in a given oxide material should be made to feel a responsibility to the scientific community

to produce a quantity of the best possible material.

A striking example in the metals field, is the contribution that the American Smelting and Refining

Company made some years ago, by producing and

making available spectroscopically pure copper.

This AS&R copper, which became the standard

research material over a period of many years, was

a major contribution to improved understanding of

the noble metals.

And now, given this improved scientific understanding of defects and transport coming from better

materials, more careful experiments, and theoretical calculation of defect parameters, how can

this information be coupled to the practical technology of ceramic materials? This question apparently cannot readily be answered, unless we intro- duce a second group of problems in the study of

oxides. Prominent among these problems are the

following: (1) the study of the role of short-circuiting

paths in diffusion and ionic conductivity, starting

from very simple dislocation arrays; (2) the adhesion

of scale during oxidation; (3) the kinetics and

mechanisms of grain growth during sintering; and

(4) the relation between chemical and tracer dif- fusion coefficients. To a purist in the field of point

defects these problems may be of lesser interest,

but, conversely, in real technological situations the

intrinsic behavior of structurally perfect crystals

may be of only minor importance. Instead, the proc- ess may be dominated by factors involved in the

above, or similar, problems. Most of these problems

received very little attention at this conference.

For problem (4) there is an elaborate formal theory

based on the thermodynamics of irreversible processes, but not much data in which the various

diffusivities are all measured on a single system.

Such a detailed study is required in order to deter- mine which simplifying approximations can be

introduced into the theory for a given diffusion

problem.

It appears that only by expanded activity in this second group of problems, in parallel with the

fundamental studies of lattice defects, can we

anticipate a useful contribution to the understanding of technologically important processes.

290-621 0-68— 15 217

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