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IMRRIR 82-2477

NATL INST. OF STAND & TECH

lllllllll aiiQb ESH75M rent Understanding of Pitting

and Crevice Corrosion and Its

Application to Test Methods for

Determining the Susceptibility to

Such Corrosion of Nuclear Waste

Metallic Containers

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Center for Materials Science

Materials Chemistry Division

Washington, DC 20234

March 1982

DISTRIBUTION OF THIS DOCUMENT IS UNLIMITED

Issued April 1 982

Prepared for

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JDepartment of Energy

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NATIONAL BURSAT

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CURRENT UNDERSTANDING OF PITTING

AND CREVICE CORROSION AND ITS

APPLICATION TO TEST METHODS FOR

DETERMINING THE SUSCEPTIBILITY TO

SUCH CORROSION OF NUCLEAR WASTE

METALLIC CONTAINERS

MAY 14 1982

/ - a :

V -

Jerome Kruger and V. Kay Hardman

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Center for Materials Science

Materials Chemistry Division

Washington, DC 20234

March 1982

DISTRIBUTION OF THIS DOCUMENT IS UNLIMITED

Prepared for U.S. Department of Energy

Washington, DC 20585

U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler. Director
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CURRENT UNDERSTANDING OF PITTING AND CREVICE CORROSION

AND ITS APPLICATION TO TEST METHODS FOR DETERMINING THE SUSCEPTIBILITY

TO SUCH CORROSION OF NUCLEAR WASTE METALLIC CONTAINERS

by

Jerome Kruger and V. Kay Hardman

Center for Materials Science

National Bureau of Standards

Washington, D.C. 20234
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TABLE OF CONTENTS

Page

Current Understanding of Pitting and Crevice Corrosion 2

Pit and Crevice Corrosion Initiation 2

1. Theories of Initiation or Breakdown 3

a. Adsorbed Ion Displacement Models 3

b. Ion Migration or Penetration Models 5

c. Breakdown- Repair 7

2. Comparison of Theoretical Initiation Models with

Experimental Results 9

a. Kinetics of Initiation--The Induction Period .... 10

b. Critical Potential for Breakdown 13

c. Alloy Composition, Structure, and Environmental

Factors 15

Pit and Crevice Propagation 17

1. Theoretical Models of Pit and Crevice Corrosion

Propagation 18

a. Theories of Pit Propagation 18

b. Theories of Crevice Corrosion Propagation 22

2. Factors and Processes Involved in Pit and Crevice

Corrosion and Their Implications for Test Method

Development 24

a. Alloy Composition and Passive Characteristics .... 25

b. Electrochemical Reactions . 25

c. Bulk Solution Composition 25

d. Bulk Solution Temperature 25

Pit and Crevice Environment 25

Testing Techniques 26

Tests to Evaluate Susceptibility to Pitting and Crevice Corrosion . 27

1. Ferric Chloride Solution—ASTM G48-76 27

2. Electrochemical Techniques 28

a. Potentiodynamic Techniques 28

b. Potentiostatic Techniques 31

c. Gal vanostatic Techniques 32

3. Modified Electrochemical Techniques 33

a. Corrosion Behavior Diagram Technique 33

b. Scratch Method 33
TABLE OF CONTENTS

Page

c. Tribo-ellipsometry 34

d. Pit Propagation Rate (PPR) Curves 35

e. Electrochemical Noise Measurements 36

f. Rotating Ring-Disk Electrode 36

Significance of Current Understanding to Testing--

Concluding Remarks 37

1. The Existence of an Induction Period 37

2. The Existence of Critical Potentials 38

3. The Role of Repassivation 38

4. The Role of the Local Environment 39

References 40

Figure Captions 46

Figures 49
Current Understanding of Pitting and Crevice Corrosion

and Its Application to Test Methods for Determining the Susceptibility

to Such Corrosion of Nuclear Waste Metallic Containers

by

Jerome Kruger and V. Kay Hardman

Center for Materials Science

National Bureau of Standards

Washington, D.C. 20234

The successful use of metallic containers for nuclear waste in a repository depends on these containers being able to resist corrosive

attack by the repository environment for a long period of time (usually set

at greater than 1000 years for the whole containment system). The most

damaging and prevalent type of corrosive attack that is most likely to

breach the container wall is localized corrosion [1]. The main forms of

this kind of attack of importance to nuclear waste problems are pitting,

crevice corrosion, stress corrosion cracking and intergranular corrosion.

This review will only consider pitting and crevice corrosion with the

latter two failure modes being considered in a future review. Pitting

corrosion is a form of localized attack that produces a penetration into a metal object at highly restricted sites. Usually these sites, called pits,

are covered with corrosion products that restrict the flow of ionic species

into or out of the pit. Thus, a pit eventually becomes a small crevice --

a crevice being formed by covering part of a metal surface by material that

restricts the access of ionic species and oxygen to that surface. Crevice

corrosion occurs when such a covered surface is placed in a corrosive

environment. The restricted environments and surfaces inside pits or crevices have been called by B. F. Brown "occluded cells" [2]. In an engineered barrier system for managing nuclear waste, it is highly likely

that the packing of the backfill against the wall of a canister will put

most of that canister's surface in a crevice and produce a very large

occl uded cel 1

In order to assess the susceptibility of various nuclear waste container

candidate alloys to pitting and crevice corrosion, it is necessary to use currently available or develop new test methods that can reliably determine
the rate of localized attack. Effective test methods for assessing susceptibility to pitting or crevice corrosion are based on a better understanding

of the mechanisms of the processes that control these two similar types of

corrosion. This review looks at our current understanding of such mechanisms

and describes the presently available test methods used to assess pitting

and crevice corrosion susceptibility. Some of those test methods have been

developed using our current understanding of localized corrosion mechanisms

as their basis.

Current Understanding of Pitting and Crevice Corrosion

There have been a number of reviews of theories and mechanisms of

pitting and crevice corrosion [3-7]. This discussion will break the subject

into two parts -- initiation and propagation processes. It is especially

important in the context of assessing the susceptibility of nuclear waste

containers to localized attack that the initiation and propagation steps be

separated because the proper choice of a test method will depend on which

step, initiation or propagation, is rate determining. For example, in a hypothetical situation if the initiation of crevice attack takes a long

time, e.g. , 50 years, the propagation rate must be slow enough to prevent

penetration for as long as 1000 years. Thus, a test that only looks at

initiation or only at propagation, would give only incomplete information.

Therefore, initiation (or breakdown) and propagation processes will be

considered separately because their mechanisms probably differ and because

their separation is necessary to develop reliable and effective test methods.

Pit and Crevice Corrosion Initiation

This discussion of initiation will make extensive use of a 1976 review

by Kruger [6] by bringing it up to date. Initiation of pitting or crevice

corrosion involves the chemical breakdown of a protective or passive film

that must be present on a metal or alloy surface in order for it to be used

effectively in a corrosive environment. Chemical breakdown (the initiation

of pitting or crevice corrosion) can be defined as those processes brought

about by chemical alteration of the film or the environment that produce

changes in the passive film so that it can no longer effectively prevent

damaging anodic currents from flowing [5]. Another way to destroy the

effectiveness of the passive film is by mechanical disruption. Mechanical

breakdown occurs by the introduction of externally applied stresses; it

will be described in a future review.

2
1. Theories of Initiation or Breakdown

A number of theories or models have been proposed to describe the

events leading to pit or crevice corrosion initiation. Most of these were

listed by Hoar [8]. In order to be successful, a model must, of course,

explain the phenomenology of breakdown. While perhaps not all inclusive

and in some cases controversial, the following phenomena are usually considered

to be the required characteristics of chemical breakdown which lead to

localized attack: (a) a certain critical potential Ep must be exceeded,

called by some the critical potential for pitting [9, 10]*; (b) damaging

species (examples are Cl or the higher atomic weight halides) are needed

to initiate and propagate breakdown; (c) there exists an induction time

separating the initiation of the breakdown processes by the introduction of

conditions conducive to breakdown and the completion of the process when

pitting or crevice attack commences; and (d) breakdown occurs at highly

localized sites.

Many models for initiation have been developed that can satisfy

the above requirements. They can be divided into three groups: (a) adsorbed

ion displacement; (b) ion migration or penetration; and (c) breakdown- repair.

As Hoar [8] has pointed out, there are common elements in all of the models,

and if a true model for breakdown is ever achieved, it probably will contain

elements of all three groups. The individual groups will be discussed in

terms of how well they satisfy the phenomenological requirements,

a. Adsorbed Ion Displacement Models

Two mechanisms have been suggested under this general model.

First, there is the mechanism suggested by Kolotyrkin [13] and Uhlig [14,15].

This model considers the passive film to be an adsorbed film (probably a monolayer) of oxygen. Breakdown occurs when a more strongly adsorbing

damaging anion displaces the oxygen forming the passive film. Once the Cl

ion is adsorbed at the surface, breakdown commences because the theory surmises,

*The situation is more complex. Pourbaix [11] has pointed out that there

exists a potential which he calls the protection potential. This

potential which is more active (negative) than is the potential below

which breakdown once initiated can be stopped. Work by Pessel and Liu

[12] indicates that when Ep is determined in a way that eliminates the

time for breakdown to occur, E = E p prot

3
the bonding of the metal ions to the metal lattice is weakened. When other

anions are present, they can compete with the Cl for sites and inhibit

breakdown [14,15]. Direct evidence that this occurs has been provided by

Rozenfeld and Maksimchuck [16] using radioactive Cl . The model explains

the critical pitting potential as that potential above which Cl adsorbs on the surface under consideration. The effectiveness of ions such as Cl in

promoting breakdown is explained by their being more readily adsorbed than

oxygen in the competition for surface sites. The effect of damaging ion

concentration on induction times can perhaps be explained by this mechanism,

but such important influences on induction time as thickness of the passive

film [17] and heat treatment of the passive film [17,18] cannot be readily

accounted for by this model. Finally, this mechanism is consistent with

localized attack since adsorbed ion displacement will probably occur at

discrete sites, defects, inclusions, etc., where displacement may be more

likely.

The second model that can be considered as falling under the

general heading of an adsorbed ion displacement model is that described by

Hoar and Jacob [8]. It has, however, elements in common with the ion

migration models also. In this model, a small number (3 or 4) of halide

ions jointly adsorb on the surface of the passive film around a lattice

cation. The probability of formation of such a high energy complex is

small and thus requires a high activation energy for formation. Once

formed, however, the halide ions will readily remove the cation from the

passive film lattice as a soluble complex (this model assumes a threedimensional passive film). The film is thus made thinner at the site where

the complex first formed and the stronger anodic field at the thinned site

will rapidly pull another cation through, where it will meet more halide

ions, complex with them, and thereby enter the solution. Thus, once started,

the passive film is locally thinned and breakdown proceeds, using Hoar's

word "explosively."

This model like the Uhlig one requires a critical potential

where adsorption takes place. It also requires strongly adsorbing anions.

Because of the difficulty of forming the transitional complex, an induction

time is required, and in fact, it was developed to explain the dependence

of the induction time on Cl concentration, temperature, and potential.

Also, it is consistent with the localized nature of breakdown. This model

4
does not appear to be applicable to all systems because Ambrose and Kruger

[18] have shown that incorporation of Cl in the passive film lattice

rather than the formation of a complex on the outside of the film accelerates

breakdown for passivated iron. Hoar [19] has also found that the proper

functional dependence of induction time with temperature found for stainless

steel, which was the basis for his model, does not hold for iron. For this

metal he proposes a penetration model.

Another model that uses a transitional complex is that of

Foley [20]. He and his co-workers have pointed out that breakdown must

involve the formation of metal -anion complexes. The solubility of these

complexes controls pit initiation. If the complex is stable and insoluble,

initiation is inhibited. On the other hand, if the complex is transitory

and easily dissociated, corrosion is enhanced and pitting is initiated.

Finally, the chemical model of Heusler and Fischer [21]

which postulates that initiation is started by the displacement of oxygen

or hydroxyl ions in the passive layer by chemisorbed chloride ions at local

sites belongs under the adsorbed ion displacement heading. This displacement

reaction results in the formation of two-dimensional nuclei of a chloride

salt. Their experimental results suggest continuous growth and disappearance

of these two-dimensional chloride islands. Such growth and currentless

chloride catalyzed dissolution of two-dimensional islands can take place on successive underlying monolayers of the oxide until the metal surface is

reached, at which time the initiation stage is over. Essentially, this

model assumes that localised chemical thinning of the passive layer is

responsible for the initiation of pitting.

b. Ion Migration or Penetration Models

The theoretical models that can be grouped under this heading

all require the penetration of damaging anions through the passive film,

the breakdown processing being complete when the anion reaches the metal-film

interface. Most of these models consider the passive film to be threedimensional. The models differ widely, however, with respect to the mode

of penetration proposed.

At one extreme is a model assuming the existence of pores in

the passive film. This model is quite old and was suggested by Evans [22]

and co-workers. More recently, Richardson and Wood [23,24] have revised this

pore mechanism and conclude that breakdown of the passive film does not occur

5
at all because there always exist pores or defects in the film that allow

instantaneous penetration and hence instant localized corrosion. One

consequence of this is that there is no induction time for the breakdown

processes to take place because there are no breakdown processes. Richardson

and Wood account for incubation by asserting that during the beginning

stages of pitting, the currents are so low at the small bare metal sites

that it appears the pitting has not been initiated, although it actually

has been going on from the instant that the surface was exposed to damaging

anions. In this way, it satisfies the phenomenological requirement for the

existence of an induction period. The other phenomena this model must

contain are also satisfied. Thus, since the pores, bare or "weak spots,"

are discrete sites, attack will be localized. The critical potential is

the potential for adsorption at the bare metal at the base of the pores,

and the role of the damaging anion is the same, and as mysterious, as it is

for the other models.

At the other extreme of the types of penetration models are those models involving migration of the damaging anion through a lattice,

via defects, or via some sort of ion exchange process. The processes which

involve ion migration in a lattice can occur in a variety of ways. Hoar,

et al [25] suggest anion entry without exchange to produce "contaminated"

passive films. The site where this occurs becomes a path of high conductivity

where high cation currents and hence the initiation of pitting are possible.

Pryor and co-workers [26,27] also propose the production of an anion contaminated

film but suggest that cation vacancies are created which lower ionic conductivity

and promote breakdown. Another ion migration scheme can involve exchange

of 0~ or OH with the damaging anion [17]. Equation (1) shows that such a process can also create anion vacancies that will further enhance the

migration of damaging anions to the film-metal interface.

solution

^ °0 ^

^®*^lattice lattice

+ 2Dnu" + 20H , . . OH solution ( 1 )

The anion lattice migration models have no difficulty in

satisfying the phenomenology of breakdown. The critical potential is, as with all the models, the potential required for adsorption, although in

6
this case, it is on the outer interface of the passive film. The induction

time is the time required for the migration of the damaging anion through

the passive film, and since this migration may be via discrete defects, the

localized nature of breakdown would be expected. The special role of

certain damaging anions is not accounted for any better than it is by the

other models thus far considered.

c. Breakdown- Repair

This group of theoretical models which involves chemically

induced mechanical disruption of the passive film is perhaps the newest.

One version was briefly described by Hoar [8]. He postulated that the

adsorption of a damaging anion on the surface of the passive film lowers

the interfacial surface tension (surface free energy) at the solution

interface until "peptization" occurs. This "peptization" results from the

mutual repulsion of the adsorbed anions. When the repulsive forces are sufficient, the passive film cracks and damaging anions can attack the

metal exposed unless conditions exist where rapid repair of the film can occur (repassivation).

Sato [28] has extended this concept of chemical -mechanical

breakdown by providing a more detailed examination of the model. He has

suggested that high fields could lead to mechanical rupture of thin films

by electrostriction pressures exceeding the compressive fracture strengths

of the film. The so-called film pressure this generates is given by:

p-p

o s(e-l)E^ _ X

8n L

( 2 )

where p = film pressure, p^ = atmospheric pressure, e = film dielectric

constant, E = electric field, y = surface tension, and L = film thickness.

It is possible that those ions considered to be "aggressive"

in promoting localized corrosion attack increase film pressure by decreasing

the surface tension term in Equation (2) through adsorption. A recent

paper by Yahalom and Pozvansky [29] gives results that indicate that this

is so. They found that halides decrease the solid/liquid surface energy.

Ep decreases linearly with log and the addition of a surface active

agent also decreases E^. They suggest that lowering the solid/liquid

7
surface energy lowers the stress for the brittle fracture of the passive

film. Unlike the other theories we have considered, this theory does not

consider the critical potential for breakdown to be related primarily to

adsorption but instead to be the "critical potential above which the film

pressure exceeds the critical compressive stress of film breakdown" [28].

Also, this model differs from the other models described in that, although

as with the others, the special role of aggressive anions such as chloride

is not clear, this model goes further and would appear to rule out any

special role at all. This is so because zero charge potentials are similar

on mercury and probably are likewise similar on oxides as well [30]. The

existence of an induction period is also difficult to explain using the

mechanical breakdown model. Finally, the localized nature of attack is

consistent with the model if one mode of breakdown involves rupture at

discrete sites as Sato does suggest. Sato points out in his description of

his mechanical breakdown theory that whether a pit nucleates or dies depends

not only on breakdown but also on film reformation. This concept is, of

course, the basis upon which all models in the group are based. Videm [31]

first proposed that it is the interplay between the breakdown (although he

does not specifically say so, this could be mechanical as Sato has suggested)

and repassivation that determines whether a passive film will suffer permanent

breakdown that leads to pitting. Videm proposed that the critical pitting

potential is determined by the relative rates of breakdown and repassivation.

Two other models rely on a dynamic breakdown- repair mechanism.

First, there is the localized acidification model proposed by Galvele [3].

He assumes that a transient breaking of the passive layer occurs through

mechanical or electrochemical means; hydrolysis then occurs during the

reaction between the exposed metal and the environment. This reaction

produces localized acidification at the site of the cracked film. The

amount of acidification necessary to prevent repassivation varies for each

metal or alloy. If the acidification is sufficient to prevent repassivation,

pitting is initiated. In this model the induction time is the time needed

to achieve a critical pH in a broken passive film. The pitting potential

is that potential needed to cause the current flowing to the initiating pit

to reach a value where a critical pH is reached that renders repassivation

impossible.

8
The last model using a breakdown- repair mechanism is the

thermodynamic model of Vetter and Strehblow [32]. They postulate a certain

relation between pit nucleation and the disorder of the passive layer which

may be of mechanical origin, for example, stresses. These stresses lead to

frequent rupture of the passive film. In nonaggressive environments the

breaks in the film are rapidly repaired. In the presence of aggressive

species such as chloride ions, pit initiation occurs when the repair process

results in the formation of a nonprotective salt at the site where the film

breaks. Such a salt forms at a potential determined by thermodynamics.

Below this potential, presumably E , thermodynamics favor oxide formation

and pitting is not initiated.

In the Vetter and Strehblow model and in all of the breakdownrepair models, the special role of the aggressive anions is that they

influence the rate of repassivation. Recent direct measurements of repassivation [33,34] have verified this. The critical potential can influence a number of factors, adsorption of the aggressive anion as in other mechanisms,

the rate of repassivation, the formation of nonprotective layers, and the

mechanical properties of the film.

The mechanism supports the existence of an induction period,

usually to build up critical acidity and/or concentrations of aggressive

anions. Finally, because attack occurs at rupture sites on the passive

film, the initiation of localized attack would be expected to occur at

discrete sites.

2. Comparison of Theoretical Initiation Models with Experimental

Results

It is evident from the discussion of the various theoretical

models that even some of the widely different models have elements in common. For example, virtually all of them require the existence of an adsorption step. There is, however, a high probability that no one model

is applicable to all systems. Thus, it will be very difficult to match

experimental results to models in an orderly way. Instead, our approach

will be to select only a few results described in the literature to emphasize

contrasting points of view. It is anticipated that by doing so, a clearer

definition of the problems will be achieved and experimental work will be

stimulated to provide new insights which may result in more effective test

methods. For a more comprehensive comparison of experimental results and

9
initiation models, see the recent review by Jani k-Czachor , Wood and Thompson

[35], Brit. Corros. J. , 1_5, 154 (1980).

a. Kinetics of Initiation -- The Induction Period

The only kinetics that need to be considered in chemical

breakdown, as defined earlier, are the kinetics of the processes occurring

during the induction or incubation period that takes place prior to the

significant increase in current that marks the onset of pitting or crevice

attack. A number of conflicting ideas, experiments, and explanations of

these experiments exist to explain this induction period. Three representative

groups will be described. A representative from the group of experiments

such as those of Uhlig [14,15] and Rozenfeld [16] which use the competitive

adsorption model for interpretation is not included because no induction

periods were reported.

Breakdown-Repair . A very extensive study of the breakdown

on pure A1 has been made by Videm [31]. His main conclusions were as follows;

(a) He found, as Bogar and Foley [36] had also shown for an

A1 alloy, that above E , there exists an incubation period before pitting

_ P

starts, when Cl is added to solution containing a filmed A1 surface. The

induction time decreased exponentially with potential, and a ten-fold

increase in chloride concentration reduced the time to initiate a pit by

ten-fold as wel 1

(b) Because he could not detect Cl by autoradiography,

except at sites where pitting attack had taken place, and because he could

not detect impedance changes after Cl soaking, he concluded that theories

for initiation of pitting of A1 requiring contamination of oxide films by

gradual chloride pick-up and subsequent change in conductivity are not

val id.

(c) He was able to measure regular current pulses during

anodic polarization which he ascribed to the development of local defects.

Recent studies of electrochemical noise by Bertocci and Kruger [37] using more sensitive techniques than Videm have bolstered his results by showing

that the magnitude of current fluctuation increased markedly when A1 is

above Ep and in the presence of chloride ions (fig. 1). Videm and Bertocci,

as well, found the pulses (electrochemical noise) both with and without the

presence of Cl in the experimental solution, but the nature of the pulses

10
changed markedly and abruptly when pitting was initiated at the end of the

induction period. This result led Videm to conclude that the Richardson and

Wood [23] suggestion that there is no pit initiation step because pitting

starts immediately at small flaws in the passive film, is invalid. A recent

review [35] points out, however, that the change in the nature of the current

fluctuations when pitting is initiated may bolster the Richardson and Wood

model because the pulses that are observed during the induction period indicate

flaws and the change in their nature is an artifact introduced by the potentiostat.

(d) The role of the chloride is not promoting breakdown

per se, but in preventing repair during a dynamic breakdown-repair type

of mechanism. Similar conclusions were reached for A1 by Zahavi and

Metzger [38].

These results of Videm' s thus point to a nonpenetration

dynamic breakdown and repair theory. It does not explain, however, how

a system predisposed to pit reaches the stage where there is sufficient

chloride available at the initiation sites to stop repassivation effectively.

For crevice attack where the crevice helps concentrate Cl , his results

would be understandable.* In the absence of a crevice, however, it is

difficult to see why at some definite time prior to pitting the Cl ions

should suddenly start to prevent repassivation unless they concentrate

in the film or somehow locally in the environment. Also, the localized

acidification mechanism of Galvele [3] described earlier might be operative.

Transitional Complex Formation . An experimental re- presentative of this group is that of Hoar and Jacob [8], whose model was described earlier. Their experimental data for austenitic stainless

steel showed that l/i, where t is the induction period, was an approximate

measure of the rate of breakdown. A plot of log t vs log C^-j- showed

that the breakdown rate is proportional to the n-th power of C^-|- with

n having values between 2.5 and 4.5. They found an activation energy of

60 kcal/mole and t to be dependent on potential. Their interpretation

*Wilde's [39] admonition that the kinetics of pitting corrosion and crevice

corrosion are different and that one must disentangle the two to assess susceptibility to localized attack is pertinent to this point.

11
of the values found for n was that 2.5 to 4.5 halide ions form a "transitional

complex" which once formed (because of the high activation energy, a difficult

and seldom event), will proceed to break down the film in the way described

in the section on models. Because of the high activation energy and the

requirement for 2.5 to 4.5 halide ions to form a complex. Hoar and Jacob

rule out an ion migration model for stainless steel. However, for other

systems, their model may not be valid. For example, Engel 1 and Stolica

[40] found a value of 1 for n for mild steel, and the induction time potential

independent. This could mean a penetration model where individual aggressive

anions were penetrating independently of each other. Foster and Hoar [19]

and also Ambrose and Kruger [18] propose a penetration model for the passive

film on iron. Other values of n have, however, been found. For A1 , Bogar

and Foley [36] found values of n between 4 and 11 while Videm found a value

of 2. Obviously, one must Be careful in using a simple interpretation such

as Hoar's and Jacob's to explain the meaning of n. Penetration . All of the penetration models explain in a very direct way the existence of an induction period. This period is the

time required for the damaging anion to travel through the passive film.

An experiment directed at providing a test of this idea was given by McBee

and Kruger [17] who measured the time to breakdown as a function of passive

film thickness. They found that the thicker the passive film, the longer

the time to breakdown. They were not able, however, to establish a functional

relation between thickness and breakdown time, since three functional relationships fitted the data equally well. To show further that penetration

is involved in the breakdown of the passive film on iron, Ambrose and

Kruger [18] showed that by removing part of the passive film by cathodic

reduction and then regrowing the film in the presence of Cl , one could

repassivate, but the time for breakdown was shortened. They interpreted

this as evidence for the penetration model because the regrowth step incorporated Cl ions. These ions, therefore, had a shorter distance to penetrate,

and the time for breakdown was shortened. Finally, Ambrose and Kruger

showed that annealing the passive film on iron at 65 C for one hour and

then bringing the system back to room temperature increased the induction

time over 100-fold. Therefore, they reasoned the defects in the film

whereby penetration proceeds were removed by the annealing. Similar results

for Ta were found by Vermilyea [41], but opposite results were found for

12
stainless steels by Okamoto [42]* (annealing shortened the induction time).

These results plus others [18] [42] which show that longer passive film

growth times lengthen the induction period, do not rule out the transitional

complex model because the production or reduction of defects in the film

can also occur on the film surfaces and thereby affect the adsorption of

damaging anions at that surface.

Two main experimental observations exist that create problems

for the penetration model. First, micro-breakdown events [31,44] as indicated

by small current or potential pulses (fig. 1) may continuously exist throughout

the induction period. If so, why is penetration needed? There exists

contradictory evidence of Cl entry into the film lattice. El 1 ipsometric

spectroscopy [17], impedance measurements [26,45], and autoradiography [46]

^ I" have indicated gradual Cl entry. Other work, however, using autoradiography

[31], impedance measurements [31,47], and Auger electron spectroscopy [48]

show no Cl uptake by the film.

b. Critical Potential for Breakdown

There are vast numbers of experiments and discussions in the

literature on the factors or underlying causes of the existence of a critical

potential for breakdown, Ep. A discussion of its practical significance

[27,49,11] and methods of measurement [50,12] will be given in the section

on test methods. Here we will detail briefly only those ideas or findings

that relate to its role in initiation of breakdown. An incomplete listing

of different explanations, of the origin of Ep follows:

E p

Determined by Potential of Zero Charge . Almost all of

the models require as their first step the adsorption of damaging anions.

Roughly, the potential above which such adsorption takes place is the

potential of zero charge (pzc). Thus, since the adsorbed ion displacement

models require the damaging anion to adsorb and thereby displace the ions

forming the passive film, such an explanation of E would be reasonable,

P

and the effect of such variables as the nature of the metal, concentration

of anions, and temperature all affect Ep and pzc [51,52]. The transition

*If the films on stainless steel are amorphous [43], the annealing could

have caused crystallization producing more diffusion paths. On the

other hand, if the films are crystalline, annealing could destroy diffusion

paths by grain growth and destruction of lattice vacancies.

13
complex model also could use the relation of pzc to Ep, not for a bare

metal, but for the passive film-solution interface. Likewise, for the

penetration models, Ambrose and Kruger [18] have suggested the relationship

of pzc to Ep. Because it is difficult to measure the pzc for solid

electrodes [51], especially when adsorption may coincide with the initiation

of breakdown processes, there have been no good experimental verifications

that Ep is related to the pzc. Galvele [53] has pointed out that Ep and

pzc may be related if E^ is the potential to initiate pitting, but they

are not related if Ep is the potential to maintain pitting.

A major problem with relating pzc to E is that the pzc

for many anions, both aggressive and nonaggressive alike, are similar on mercury and probably are similar on oxides as well [30], whereas the values

of Ep can vary a great deal

E Determined by Film Pressure to Fracture Film. Sato

-P ^ [28] has identified E with the critical potential above which the film

P

pressure exceeds the critical compressive stress of film mechanical

breakdown. The film pressure, p, resulting from both electrostriction

and interfacial effects is given in Eq. (2). Using the above assumption,

Sato has been able to derive the relationship between the Ep and the concen- tration of anions in solution as follows:

(3)

dlnaj^^ e(e+l)E

where a^^^ is the activity of the salt MA, r^- is the adsorption density

of A , k is the Boltzman constant, T the temperature, E the electric

field strength, and e the film dielectric constant.

He was able to calculate from Eq. (3), using the results of

Leckie and Uhlig [15] for co-existing anions, that the density of adsorption

of the aggressive Cl ion on passive stainless steel is of the same order as

- - 2

- that for the inhibitive OH , NO^, and SO^ ions. One problem with the relation

of E to film pressure is that the special role of aggressive anions such as

Cl is uncertain because they have about the same polarizability as oxygen ions

and only slight changes in dielectric strength (a few percent) might occur if

14
they were to replace the oxygen ions. This would be particularly true at

the low concentrations of Cl to be expected in passive films undergoing

breakdown [54].

Ep

is Determined by Repassivation Kinetics . This concept of

E is based on kinetic considerations. Some studies [31,55,38,42] have

P

found that a competition between repassivation and breakdown of films on A1

is important in determining initiation of final breakdown. Videm [31]

proposed that the value of E is determined by the relative rates of these

P

two processes. He based this on his experiments that showed a narrow

potential region where repair and breakdown were competitive. Below this

region, repair predominated and above it, breakdown was the major process.

Because potential may affect repassivation kinetics [33,34] and because, as Videm and others [53,38,42,55,56] suggest, initiation of localized attack

depends on a dynamic balance between breakdown and repair, the kinetic

control of Ep must definitely be considered. More quantitative details are

still lacking, however, and the problem remains of finding a way to concentrate

damaging anions before final breakdown occurs when no crevice or "occluded

cell" situation exists. Galvele's [3] local acidification model may provide

a way for this to happen.

Ep is Determined by the Potential of Formation of an Unprotective Film . One possible factor determining E which may not apply to all

P cases may be that it is the potential above which an unprotect ive film may

form. The thermodynamic model for breakdown [32] described earlier suggests

this. Ambrose and Kruger [18] found that this may be the case for iron.

They found that the E they measured in pH 8.4 borate solutions containing

Cl was both equal to the pzc for iron found by Lorenz and Fischer [57] and

the equilibrium potential for the reaction that produces in a series of

steps the nonprotective y-FeOOH [58]. It was proposed that y-FeOOH was formed after penetration by Cl had produced an ion cqnductive path for

2+ Fe which would produce the nonprotective y-FeOOH above Ep. The role

played by Ep in producing a nonprotective film in the presence of Cl is

shown dramatically in fig. 2. Whether this sort of explanation of Ep can apply to a large variety of situations remains to be shown.

c. Alloy Composition, Structure, and Environmental Factors

It is not possible within a limited space to discuss the many

complex and contradictory results concerned with the effect of alloy composition and structure and environmental factors that influence breakdown

15
initiation processes. An excellent review has been given by Smialowska

[4]. This discussion will concentrate on how some of these factors relate to

the theoretical models for breakdown.

Alloy Composition . Kolotyrkin [5] points out in his

review of pitting that the tendency to breakdown decreases with an increase in the content of N, Ni , Cr, and Mo, especially for the latter

two. The main effect of these beneficial alloying elements is to shift

Ep in a noble direction. This explains phenomenologically why they are beneficial but does not actually tell us how they work. One way in

which they can retard initiation processes is by producing a passive

film that is more difficult to penetrate because it provides fewer

diffusion paths. Hoar [58] has suggested that some of these beneficial

alloying elements can do this by producing amorphous (glassy) films.

McBee and Kruger [43] have found experimental evidence for this when Cr

is added to Fe and also have found reduced rates of penetration [59].

Another way, however, in which these alloying elements can affect breakdown

is by affecting repas si Nation kinetics. Mo is known to be beneficial

in retarding breakdown leading to pitting [60]. Measurements by Ambrose and

Kruger [55] indicate that its addition to an alloy does increase the rate of

repassivation.

Alloy Stigiijiture . Only the surface structure of alloys

will be considered here. Three kinds of possible variations in surface

structure have been studied. First, Kruger has looked at the role of

crystallographic orientation on breakdown tendency for iron [62]. As

Fig. 3 shows, the tendency to pit goes up as the surface approaches the

{100} orientation. He also found that pitting density varied with

crystallographic orientation [63]. Another example is the single

crystal work by Smialowska and colleagues [64,65] on Fe-16Cr. They

found that nucleation occurred at crystal subboundaries as well as at

other sites where there were not necessarily any metal inhomogeneities.

The second surface variation can occur by the introduction

of defects by abrasion, cold work, or other deformation processes.

*More recent work by Kodama and Ambrose [61] found that the addition

of Mo affects the propagation stage of pitting more profoundly

than the initiation stage.

16
Contradictory results have been found. Thus, Rozenfeld [66] found that

removing cold work by annealing an 18Cr-9Ni-Ti steel increased the number

of pits while Randak and Trautes [67] found that cold work increased the

number of breakdown sites for an 18Cr-8Ni steel.

Finally, surface heterogeneities play a big role in initiating

breakdown. Smialowska, Jani k-Czachor, and co-workers [65,68,69] have found

extensive evidence for the nucleation of pits at inclusions such as sulfides

or chromium oxide for a number of alloys. Bond and co-workers [70] showed

that very small amounts of segregated metallic impurities in quite pure

aluminum single crystals could produce breakdown nucleation sites.

The obvious, but sometimes contradictory, effect of surface

structure on breakdown can be explained by all the theoretical models

discussed because it is reasonable to expect adsorption, penetration,

complex formation, or repassivation kinetics to be affected by surface

heterogeneities.

Environmental Factors . These factors can include solution

composition (including pH) and temperature. There are large numbers of

experiments in the liteiri&ture [4,5] and the results are contradictory.

Generally, however, breakdown tendency increases with temperature and

concentration of damaging anions while it is not strongly dependent on pH

[71]. The nature of the^^ anions in solution play a large role [15] especially

if competitive anion adsorption situations are possible. There is some controversy as to what cpnstitutes a damaging anion. Thus, Hoar [72]

considers the CIO^ ion as damaging as Cl , whereas Uhlig [15] has found it

to be inhibitory. With regard to models, again all models will depend on environmental considerations. A detailed examination of how all the various

and contradictory results in the literature best fit the various models is

outside the scope of this discussion. For such a discussion, see references

3, 4, and 35.

Pit and Crevice Propagation

After pitting or crevice corrosion is initiated by the breakdown of a passive film on a metal or alloy surface, the damaging stage of those two

localized corrosion failure modes, propagation, commences. If the rate of

propagation is very rapid, initiation would be the stage that would control

the life of a nuclear waste container. As stated at the beginning, it is a requirement of an effective test method to determine whether initiation or
propagation is rate determining. This section describes our current understanding of the factors that control the mechanism and hence rate of propagation.

1. Theoretical Models of Pit and Crevice Corrosion Propagation

As can be inferred from our discussion of the theories of initiation,

there is not always a sharp separation between propagation and initiation

theories. There are models in the three groups of models described for

initiation or breakdown that must of necessity relate to propagation processes

as well. Thus, Kolotyrkin [5] and Rosenfeld [81] explain both stages using

adsorbed ion displacement models. The version of the ion penetration or imigration model requiring pre-existing flaws as espoused by Richardson and

Wood [24] has been characterized as being equally applicable to propagation

as it is to initiation [35]. In fact this characterization has been used

to argue that no initiation stage exists [35]. Finally, the localized

acidification [3] and the thermodynamic versions [32] of the breakdown- repair

models have difficulty in separating initiation from propagation, and their

proponents deal extensively with the propagation stage.

In this section we will emphasize the theoretical aspects or mechanisms that apply most directly to propagation of pits and crevices.

Since the geometry of crevices differs greatly from pits and crevice corrosion

starts at lower potentials [3], the theory of crevice corrosion propagation

will be treated separately.

a. Theories of Pit Propagation

The theories of pit propagation are, of necessity, less

atomistic in nature and are more concerned with geometrical and mass transfer

considerations. We will consider both theories that combine the initiation

and propagation stages and those that clearly only deal with the propagation

stage.

In 1967, Rosenfeld and Danilov [73] invoked the adsorbed ion

displacement model to explain the initiation of pits at sites, where the

oxygen-metal bonds were supposedly weaker than on most of the metal surface,

by the displacement of oxygen ions with chloride ions. This displacement

of oxygen by chloride ions leads to a penetration of the passive film at

highly localized sites. This event creates local anodes and cathodes and

the concentration of aggressive anions (e.g., chloride ions) is enhanced at

the positive anodic sites. This situation creates favorable conditions for

18
raising the potential in the newly started pits and thereby further stimulating

the anodic reactions (metal dissolution) taking place in the new pits.

These pits become covered over with corrosion products and thereby become

small crevices where Rosenfeld suggested the mechanisms for crevice corrosion

become operative. We will describe those mechanisms in our discussion of

crevice corrosion.

Galvele's localized acidification theory [3] provides a great deal more detail to the Rosenfeld picture just described. The initiation

stage described by the localized acidification theory has been discussed

earlier so we will begin in our description of Galvele's theory at the

point in time after a pit has been initiated. Galvele has used the diagram

(a simplified Pourbaix or potential -pH diagram) given in fig. 4 to describe

events that take place at the beginning of the propagation stage. In

fig. 4, a represents the pH of the bulk environment and b the pH in the

starting pit where conditions for the metal dissolution reaction

Me = Me^"*" + Ze' (4)

becomes thermodynamically favorable. The hydrolysis reaction

Me + ZH

2

O = MeCOH)^ + ZH'^ + Ze" (5)

only takes place between a^ and in fig. 4. When the potential is above

82 , the localized acidity produced by reaction (5) will promote dissolution.

Potentials between a^ and ^2 will promote repassivation. A somewhat similar

description of the role of local acidification on pit propagation has been

given by Pourbaix [9].

An important aspect of Galvele's mechanism is that it provides

insights into the effect of potential on the rate of metal dissolution in

pits. The rates of reactions (4) and (5) and the hydrogen evolution reaction

+ 2 e" = H

2

( 6 )

are potential dependent. If the hydrolysis reaction producing MeCOH)^ is

rapid, reaction (5) will give the rate of proton production, the higher the

potential the higher the rate. When the potential is low, however, reaction ( 6 )

19
is favored and protons are consumed. When the potential is at a value

where the rates of reactions (4) and (6) are equal, it is the corrosion or open-circuit potential in an acidified solution close to the anode, called

* * by Galvele. At E^, rate of the proton production reaction (5) is

equal to the rate of the proton consumption reaction (6), and no acidifica-

* tion will take place. Below E passivity conditions exist and any passive

c ^ film that is disrupted will be reformed rapidly. Above E^ repassivation is

nonexistent or ineffective, and the pit will propagate. This treatment of

*E^ by Galvele suggests a somewhat, but not completely, similar explanation

by Pourbaix of his protection potential, The various potentials

used to characterize pitting phenomena will be described more fully in the

section on test methods.

Vetter and Strehblow [32] in the exposition of their mechanism

of pit growth (and initiation) questioned the local acidification mechanism

just described. Their calculations indicated that no significant changes

in solution composition could take place during initiation or when very

small pits were growing. (A newer calculation by Galvele [74] indicates

significant changes of pH but insignificant increases of chloride ions

during the early stages of pitting.) The theory of Vetter and Strehblow

was described briefly in the discussion of initiation. Their theory is an elaboration of propagation ideas developed earlier by Franck [75].

The ideas of Vetter and Strehblow on initiation have already

been discussed. Their picture of pit growth starts with their attempts to

resolve the problem of the simultaneous existence of active pitting sites

and passive regions on the same surface. This requires a potential difference

between the surfaces in pits and the rest of the surface and indeed potential

drops of tenths of volts have been observed. However, Vetter and Strehblow

have shown experimentally [76] that such large differences cannot exist in

small pits. Based on calculations, they dismiss the cause of this potential

difference as being due to large differences in pH between the environment

in the pit and that outside, as suggested by Galvele [3] and others.

Instead they use Franck's proposal [75] that a highly resistive surface

layer, probably a salt, must exist on a growing pit's surface.* This salt

*Strehblow [77] recently included the possibility that local acidification

due to hydrolysis may occur in the pit along with the formation of a salt

film especially in cases where a dense film may not form, e.g. , for aluminum.

20
layer was said by Vetter and Strehblow to be formed by the cations of the

metal and the aggressive anions. They hypothesized that it is dense,

poreless and contiguous to the pit solution which is not saturated with the

salt. The current flows through this layer in the same manner as the

current has been proposed to flow through an oxide film, as an ionic current

produced by a high electric field (''' lO^V/cm). A key point in their mechanism

is that the rate of pit growth (corrosion) is equivalent to the rate of

salt film dissolution. At steady state the dissolved metal ions are replaced

by the anodic ionic current which is equal to the steady state corrosion

rate (corrosion current density). Since the thickness of the salt layer

determines the field strength which sets the ionic current, it is regulated

by the process to set the steady-state value of the corrosion current

density. Since many workers have found pits with hemispherical shapes

[75-79], the corrosion current density must be constant over the whole

surface of such pits, indicating a very uniform salt layer in the pit.

More recent ideas by Strehblow [77] indicate that since repassivation of

pits can take place when the potential is lowered below that repassivation is closely related to the removal of accumulated aggressive anions

into the bulk electrolyte, a diffusion-controlled process. This bolsters the

idea that aggressive anions in the pits are essential for stable pit growth.

The final model type for pit growth to be described is one that considers diffusion-controlled or mass-transfer processes as being of great importance. Kaesche [78] in considering the pitting of

aluminum in a NaCl solution used a mass transport model. He assumed the

following: (a) the base of the pit was active and (b) hydrolysis of the

3+

A1 ion lowered the pH in the pit which caused a buildup of an AlCl^

layer at the base of the pit because of the dissolution of the AKOH)^

formed during hydrolysis. Using these assumptions, he combined a one- dimensional solution to the mass transport equations for A1 , Cl and

Na^ ions with the condition of electroneutrality to determine the

limiting current density for aluminum ion dissolution at the base of a 3+

pit. He calculated that the limiting current density, ij^, for A1 in a single pit was

21
. _ ZFKD

\ 6

( 7 )

3+ where Z = 3 (for A1 ), F is the Faraday constant, D is the diffusion

3+

constant for A1 , 6 is the thickness of the diffusion layer and K is a constant that has the dimensions of concentration and is determined from

model equations as a function of bulk concentrations of NaCl

Beck [79] modified this model in order to treat the pitting

of titanium because the Kaesche model did not reflect the phenomenology of

titanium pitting. Beck spelled out the following assumptions for his model

for titanium;

(a) Base of pit is active.

(b) Titanium tetrahalide (TiX^) is formed at base of pit.

(c) Covalent TiX. diffuses away from the surface.

^ - +

(d) TiX^ is hydrolyzed in solution to Ti02, X , and H by

first-order reaction with rate constant,

k (sec ^).

(e) Bulk solution has concentration, C° of HX.

(f) Rate of dissolution of Ti is limited by mass transport

of X to base of pit.

(g) One-dimensional mass transport for TiX^, X , and

through diffusion layer of thickness 6.

(h) ElectroneutVal ity.

A mathematical treatment using these assumptions lead to an expression similar to Eq. (7) but.with K replaced by a factor R that depends

on the regeneration of halide ions by hydrolysis within the diffusion layer.

R is the ratio of the experimentally measured current density at a particular

halide concentration to that calculated from the simple expression for the

limiting current density. The model predicts that if hydrolysis is too rapid

a passive oxide layer would form too close to the pit's metal surface and block

the surface. Beck suggests that this will happen for iodide. For bromide the

hydrolysis rate will be in the right range and it will be too slow for chloride

sol utions.

b. Theories of Crevice Corrosion Propagation

The propagation of crevice corrosion differs from pitting

corrosion mainly because of two reasons: (a) the geometry of a crevice

22
is usually different from that of a pit, being much flatter and shallower

and (b) the restriction of mass transfer of reactants is much more effective for a crevice than for a pit. In an admirable and quite

useful effort to develop a mathematical model for crevice corrosion,

Oldfield and Sutton [81] have described rather completely the mechanisms

of crevice corrosion propagation. To do this they have incorporated the

earlier mechanistic ideas of Rosenfeld [81], Karl berg and Wrangler [8]

and Fontana and Greene [83].

Oldfield and Sutton divide the processes controlling

crevice corrosion into four stages with the first two stages being

involved essentially in the initiation of breakdown of passivity and the

latter two being involved in the actual removal of metal within the

crevice. Their four stages are as follows:

Stage 1 . During this stage the normal corrosion reactions

that take place on an uncreviced surface take place outside and inside a crevice. Assuming a NaCl solution containing dissolved oxygen, the

reactions at the anodes and cathodes of a metal surface bearing passive

film wi 1 1 be

Anodic reaction: Me = Me^^ + Ze

Cathodic reaction: H

2

O + 2e + 1/2 O

2 = 20H

The overall reaction when these half reactions are combined

2Me + I

^ 2Me(0H)2 (10)

leads to the thickening of the passive film. If the diffusion of oxygen

into the crevice is severely limited by a deep tight crevice, the crevice

solution eventually becomes depleted in oxygen because of reaction (9),

and the first stage is at an end.

Stage 2 . After the depletion of the oxygen in the crevice

solution, the anodes and the cathodes of the metal surface become localized

with the cathodic reaction taking place outside of the crevice and the

crevice reaction inside. The production of an increase in the concentration

of metal ions by the anodic reaction inside the crevice and their participation

in a hydrolysis reaction (Me^^ + ZH

2

O = MeCOH)^ + ZH^) causes the precipitation

of metal hydroxides and a lowering of the pH. In order for electroneutrality

( 8 )

(9)

23
to be maintained and because of the mass transport associated with carrying

the current during the reduction in pH, cations move out of and anions move

into the crevice. In this way the concentration of chloride ions is built

up in the crevice solution thereby increasing its corrosivity. The rate of

this increase in crevice solution corrosivity depends on the magnitude of

the corrosion current and the volume of the crevice (and hence its average

gap).

Stage 3 . When the concentration of Cl and the pH of the

crevice solution reach values where repassivation becomes slow or not

possible, stage 3 commences. This stage is the point in the process where

there is a complete breakdown of passivity. The crevice solution composition

in which this occurs was called by Oldfield and Sutton the "critical crevice

solution" (CCS). This solution can be defined by its pH and chloride

concentration. They stated that the value of the passive current (the

corrosion current before the breakdown of passivity) and the CCS are factors

that are important in determining an alloy's resistance to crevice corrosion.

Stage 4 . During this stage, which begins with the complete

breakdown of passivity within the crevice, propagation, i.e., rapid dissolution

of the alloy inside the crevice, takes place. Outside the crevice the

metal dissolution reaction is, of course, balanced by rapid oxygen reduction.

There may also be some hydrogen evolution occurring in the highly acidic

solution inside the crevice.

Using these four stages, Oldfield and Sutton have built a mathematical model that emphasizes the importance of the CCS and the passive

current and use these factors to characterize an alloy's resistance to

crevice attack, crevice geometry and the chloride concentration in a bulk.

Using the model they are able to predict whether the pH will reach a limiting

value controlled by mass transport or go beyond the limiting value to a critical value where breakdown and rapid corrosion takes place. A complete

discussion of their model is beyond the scope of this review.

2. Factors and Processes Involved in Pit and Crevice Corrosion and Their

Implications for Test Method Development

A brief listing of the factors as suggested by the theories and

mechanisms discussed above will now be given. A good starting point is

fig. 5 from the paper by Oldfield and Sutton [80]. It is given for crevice

corrosion only but, with the exception of crevice-type and crevice geometry.

24
the same factors apply to both forms of localized corrosion. Our listing

will not include all of the items in fig. 5, but it will add some not in

that figure.

a. Alloy Composition and Passive Film Characteristics

As mentioned in the section on pit and crevice initiation,

the constituents of an alloy and especially their effect on the nature of a passive film are of crucial importance. Test methods used for a given

alloy must assure that the test environment has a composition that includes

major and minor constituent ions that may be present in the crevice or pit

solutions.

b. Electrochemical Reactions

All electrochemical reactions, anodic and cathodic, should

be considered in test designs. In crevice corrosion situations, cathodic

reactions may play a bigger role than hitherto expected. Another group of

electrochemical reactions that is usually neglected because it is of a transient nature is repassivation. Because many models emphasize the

importance of repassivation in pitting or crevice attack, test methods that measure it are of great value.

c. Bulk Solution Composition

Many of the models indicate that except for the effect of

the bulk solution on the pit or crevice environment, the bulk environment

is of secondary importance. This is especially so if a test environment is

sought that simulates propagation processes.

d. Bulk Solution Temperature

In the corrosion of containers for nuclear waste containers,

temperature is perhaps the most important factor here. It affects passive

film characteristics, the kinetics of corrosion, the rate of repassivation

and the composition of the pit or crevice environment. It should always be

a parameter in a test program that seeks to predict long-term susceptibility.

Pit and Crevice Environment

As stated above, most pitting and crevice corrosion mechanisms involve

a change in the local environment. It is of crucial importance to characterize

this environment for a given bulk environment-alloy system and to determine

whether this local environment is the best test environment to use in the

development of an accelerated test method.

25
Testing Techniques

Tests designed to determine the susceptibility of alloys to localized

corrosion, specifically pitting and crevice corrosion, often include natural

exposure tests. These tests involve placing metal coupons or assemblies

(see fig. 6b and c) in the environment, i.e., seawater, and either monitoring

the sample ijn situ or leaving it there over a long period of time (greater

than one year). After the samples are retrieved, they are weighed and

examined. The weight lost of the sample is then used to determine the

corrosion rate. The Multiple Crevice Assembly with 40 small crevices on

it, developed by Anderson [84], is generally used to study crevices. The

large number of crevices is desirable in a statistical approach to the

results.

Natural exposure tests are impractical for metallic containers in

nuclear waste repositories in that information is needed about the localized

corrosion resistance of alloys over 1000 years. In this situation accelerated

testing technique, capable of determining pitting and crevice corrosion

rates of selected alloys at various temperatures, are needed. Corrosion

rates, as previously mentioned, are dependent on both initiation time and

propagation rates of pits and/or crevices; therefore, techniques that can separate the two are desirable. Unfortunately, most techniques involve

both, but if the experiments are carefully designed this conflict can be

greatly reduced or eliminated.

Because of the requirement that nuclear waste containers resist localized

corrosion for very long periods, evaluations of their susceptibility must,

of necessity, be carried out using accelerated tests. This discussion will

only deal with such tests. While most accelerated testing techniques

cannot produce absolute numbers for corrosion rates of a particular alloy,

these methods can rank alloys as to their susceptibility to localized

corrosion and can describe mechanisms involved with initiation and propagation

of the pit and/or crevice [85]. To evaluate the suitability of alloys for

application as containers of nuclear waste, rates for pit or crevice initiation

and propagation are needed and accelerated tests that only rank alloys are not enough.

26
Tests to Evaluate Susceptibility to Pitting and Crevice Corrosion

1. Ferric Chloride Solution - ASTM G 48-76

This is the standard test method for crevice corrosion resistance

of stainless steels and related alloys. The ferric chloride solution seeks

to simulate crudely the pit or crevice environment. The crevice design

consists of two TFE-fluorocarbon blocks on either side of the sample and is

held together with two rubber bands at right angles to each other. Metal

coupons (for pits) and crevice samples are immersed in 10 percent FeCl^ at

a controlled temperature [86]. The FeCl^ solution is a highly corrosive

environment similar to that found in the pit environment which accelerates

the corrosion rate of alloys susceptible to localized corrosion. Samples

are weighed and optically examined for pits and crevices. Those alloys

that have low corrosion rates and exhibit little or no evidence of pitting

are considered resistant to localized corrosion.

Another approach or test method to determine crevice corrosion

susceptibilities is the use of the crevice corrosion temperature by Garner

[86]. In this test method the critical temperatures of various alloys are determined using FeCl^ solutions. After the crevice assembly has been

immersed in 10 percent FeCl^, the temperature is increased by 2.5 C every

24 hours from -2.5 C up to a temperature where crevice corrosion can be

observed. This temperature is called the crevice corrosion temperature

(CCT). Garner's results showed an increase of the CCT of stainless steels

with the addition of 6 percent Mo and that these metals would be essentially

free of crevice corrosion in ambient sea water.

The ferric chloride solution is not an ideal accelerated test for

several reasons. Since ferric ions are always present in the crevice, the

cathodic reaction for the corrosion process can occur. This may suggest

possible susceptibility of stainless steel and iron alloys but is an unnatural

crevice environment and therefore of little value for other alloy materials

such as those of titanium. It is also difficult to reproduce the exact

crevice geometry from test to test. Even for stainless steel tests, the

ferric chloride test represents propagation conditions only and ignores the

initiation process necessary to develop the cathodic reduction of oxygen

outside the crevice.

27
2. Electrochemical Techniques

Polarization curves are obtained to (a) acquire general information

on the localized corrosion process, (b) determine critical potential values

such as critical pitting potential, critical crevice potential, protection

potential, and the corrosion potential, and (c) suggest the composition of

the local environment conditions such as pH and chloride content. The

effect of chloride and pH content can be seen by comparison of the various

concentrations on the polarization curves. Three methods are most often

used in these studies and will be discussed separately.

a. Potentiodynamic Techniques

Potentiodynamic anodic polarization of an alloy gives the

current-potential relationship for the alloy in a given environment by

scanning at a selected rate over a range of potentials. When an alloy

sample is immersed in an aqueous .solution, oxidation and reduction occur on

its surface. Although the alloy can function as a cathode or an anode,

corrosion of the alloy generally results from anodic currents. The corrosion

potential, or open circuit potential, is the potential, relative to a reference electrode, at which the rate of oxidation is equal to the rate of

reduction. In other words, the magnitude of the anodic current is equal to

the cathodic current. The corrosion potential (-163 mV) is shown in fig. 7

at A. Potentials less than (lower than A in fig. 7) are in an active

region, and potentials greater than considered increasingly noble

potentials. The general corrosion rate can be determined from this by

extrapolating a straight line through the cathodic portion of the curve (in

the active region) to E . That intersection defines the corrosion rate

of the alloy and is 6.7 x lO'^ na/cra for 304 stainless steel in this example.

The corrosion rate just mentioned does not consider localized

corrosion and, therefore, is not particularly useful if pitting or sites

for crevice corrosion develop. Under localized corrosion conditions, the

corrosion rate of the alloy will increase greatly at highly localized

sites. Potentiodynamic anodic polarization scans can give qualitative

information of the alloy such as its passivation tendencies and effects of

inhibitors, under different environmental conditions. Critical potentials

and potential regions of the scan, however, can only indicate the relative

susceptibility of an alloy to pitting and/or crevice formation because

scanning over a potential range does not always allow for induction times

28
to be exceeded so that pitting is initiated at the correct Ep. Also, the

potential region of the alloy's passive film and its current density can be

roughly determined.

A potentiodynamic anodic polarization curve is shown in fig. 7

from A to B. The potential of the sample was scanned slowly in the

anodic direction from the corrosion potential to B. Pits and crevices are initiated at potentials more noble than the critical pit or crevice potential,

Ep (see fig. 7). This potential range is called the pitting region. At

potentials above Ep^^^ the alloy is susceptible to pitting. At potentials

lower than the protection potential, Ep^^^, pits and crevices are repassivated.

At potentials lower than the protection potential, Ep^^^, pits and crevices

are repassivated. A large passive region, between A and Ep, is also desirable

in that it indicates the oxide film is stable over a wide potential range.

Pitting and crevice corrosion can occur simultaneously unless the alloy

sample is carefully prepared. If the critical crevice corrosion potential

is desired, then a crevice assembly is needed (see fig. 6d and e). However,

crevices should be avoided in pitting experiments (fig. 6a).

Defranoux [90] and Wilde and Williams [49,91] have studied

pit initiation on stainless steel and nickel alloys in chemical and marine

environments. Specifically, they were comparing results of accelerated

testing by polarization techniques to natural exposure (weight-loss) tests

and found that accelerated testing can represent the relative susceptibilities

of different alloys to pit corrosion. However, these results were meaningless

if crevices were present. Their observations suggested that the propagation

mechanisms of both pitting and crevice attack are identical but the initiation

mechanisms are different. Since the initiation in crevice corrosion was found to have faster kinetics [92], the critical pitting potential can be

related to the breakdown of passivity by pit initiation only when no crevices

are present. Unlike pit initiation, crevice initiation [93] results from

differential aeration or a depletion of the cathodic reactant which fosters

the breakdown of the passivity in the crevice.

Pit and crevice propagation, however, have been observed to

follow the same mechanism. It involves increasing anodic dissolution in

regions of limited diffusion. Therefore, the experimental designs of differentiating pits from crevices are not as critical when studying propagation

effects of localized corrosion. One method of determining the relative

29
amount of propagation of pits and/or crevices in alloys is by comparing the area in the hysteresis loop formed during cyclic or pitting polarization

curves. Cyclic or pitting scans can be obtained by reversing the

potential sweep in the cathodic direction at a particular potential or current (B in fig. 7). As the pits repassivate during the reverse scan,

the current is reduced and a hysteresis loop emerges. The protection

potential, is determined by the intersection of the forward and reverse scan. Pitting characteristics of the alloy are represented by the

shape and position of the hysteresis loop. A rule of thumb is that the

difference potential, Ep - Ep^^^, is related to the crevice corrosion

resistance of the alloy and is used to represent the area of the hysteresis

loop. In fig. 7 the sample was severely pitted before the potential was

reversed. Not surprisingly, the material did not repassivate quickly and

indicated high susceptibility to crevice corrosion under these severe

conditions.

An example of the application of cyclic scans to crevice

corrosion susceptibility can be found in the study of crevice corrosion of

Hasten oy C, Incoloy 825, and Carpenter 20Cb3, combining the polarization

measurements with weight- los^^ data [49]. An artificial crevice assembly

was used in the tests, and the alloys were subjected to potentials large

enough to break down the passivity through crevice initiation. It was

found that under increasing anodic conditions, premature breakdown of

passivity occurred when a crevice was present. In the absence of a crevice,

no breakdown of the passivity would occur (no hysteresis). The cyclic

polarization results can be seen in fig. 8 and agree well with crevice

corrosion weight-loss data of these alloys exposed to seawater for two

years. The weight loss for Hastelloy C, Incoloy 825, and Carpenter 20Cb3

2

is 0.16, 4.1, and 26.1 mg/cm , respectively. The cyclic potent! odynamic

anodic polarization curve for Hastelloy C exhibits no hysteresis loop (note

reversibility of the curve). The curve for Incoloy 825 has a small amount

of hysteresis (very little crevice corrosion) while Carpenter 20Cb3 exhibits severe crevice corrosion.

There are several problems involved with potentiodynamic and

cyclic potentiodynamic anodic polarization curves. The presence of a

30
crevice makes it very difficult to detect pitting susceptibility. Any

pitting values for and Ep obtained in the presence of a crevice would

be suspect. This is illustrated in fig. 9 for Fe-30Cr-3Mo in nitrogen

saturated 1 M NaCl at 25 C [49]. There was no pit initiation, and the

major sites of dissolution were the grain boundaries where chromates and

molybdates were formed. A photomicrograph of the surface confirmed that no pits were present. However, the cyclic polarization of the alloy containing

a crevice assembly resulted in a pronounced hysteresis loop indicating

initiation of crevice corrosion.

Critical values such as Ep and Ep^^^ found through potentiodynamic scans are very sensitive to the conditions under which the experiments

are conducted. The critical pitting potential can differ widely depending

on the potential scan rate as seen in fig. 10 [94]. When the potential is more noble, the initiation time is shorter for pit formation. Therefore, a rapid increase of current can be measured when there is a short induction

time which corresponds to fast potential scanning rates, thereby giving too

low values for Ep. Thus, for long term situations like nuclear waste

containment a test should not ignore long induction times, e.g. one year, a small fraction of the 1000 years. The protection potential, Ep^^^, is not

a value unique to the alloy and environmental conditions but rather reflects

the experimental conditions necessary to repassivate the pits depending on the amount of pit propagation before the reverse scan begins [49]. This

can be seen in fig. 11 where the amount of pit propagation was increased

after pit initiation had occurred. There was a large difference in Ep^^^

particularly after the pit growth was close to 1 A/cm^ before reversing the

scan.

b. Potentiostatic Techniques

The potentiostatic technique can be used in studies of

either initiation or propagation mechanisms and involves monitoring the

current or current density versus time at a constant applied potential

[50,95]. Often potentiodynamic or cyclic polarization scans are taken

first to determine Ep and Ep^^^ values. For comparison a cyclic polarization

curve of the potential versus the log of the current is shown in fig. 12(a)

with the critical pitting and protection potentials labeled. Standard

mathematical practice is to plot the dependent variable (in this case log

i) as the ordinate with the independent variable (the potential) as the

31
abscissa. However, ASTM specification G-3 calls for data to be presented

in the reverse way with the potential as the ordinate and the log i as the

abscissa (fig. 7). This allows the longer axis of most x-y recorders to be

used for recording the current or current density.

There are two variations of potentiostatic measurements that

are commonly used. One, primarily associated with initiation studies,

measures the change in current density over time at potentials greater than

and less than the critical pitting potential, E . A new sample is used

P

with each curve. This variation is illustrated in fig. 12(c) in which E

2

is equal to Ep. When E < Ep as with the potential E-j , the current decreases

with time since the alloy's surface is passive. At potentials greater than

E , pitting is initiated; thus the current increases. Propagation studies

involve activating the alloy's surface to form pits by applying a potential

greater than E , then quickly shifted it to values less than E . The

results of this procedure can be seen in Fig. 12(d). Propagation of already

existing pits at E > Ep^^^, continues and is shown by an increase of the

current with time. However, at E < Ep^Q^> piTs are repassivated and the

current decreases with time.

The potentiostatic method of current versus time plot at

constant potential can be determined on the clean surface of the alloy.

This gives a reliable value of Ep but is very time consuming since the

induction time can be very long. Activated surfaces can, however, give

reasonably good values of Ep^^^ which some believe is the true Ep [12].

c. Gal vanostatic Techniques

The gal vanostatic technique measures potential versus time

at constant current. Generally measurements are made as the potential

decreases in time until a constant potential is reached. This potential is

the protection potential, Ep^Q^> and is shown in fig. 12(b). This technique

can take on various forms in that it can involve two dissimilar metals in a solution or the two metals may be in two different solutions which are electrically connected by a salt bridge.

Gal vanostatic measurements of potential versus time at

constant current are very quick and easy but the periodic oscillations of

the potential make interpretations of the data very difficult. Also, the

maximum potential is often unmeasurable without an oscillographic device

since the localized corrosion can occur within seconds.

32
3. Modified Electrochemical Techniques

a. Corrosion Behavior Diagram Technique

Corrosion behavior diagram techniques are useful as a qualitative

approach for comparing various alloys in a specific environment [96]. The

CBD technique is particularly valuable for determining the primary passive

potentials, E , and critical anodic current density I , in self-passivating

PP ^ samples. Both E and I values arise when the cathodic curve intersects

^ PP c the anodic curve in the active region. This is a theoretical curve since

experimentally only the net current can be measured. Self-passivation can occur in potentiodynamic scans in that an initial delay before data collection

is generally present and the scanning rate is usually 1 mV/s or less.

There is no initial delay in the CBD technique because the run begins

immediately when the sample is immersed and the scanning rate is 10 mV/s

or greater.

Figure 13 shows a CBD plot. The forward scan (scan 1) is

shown as a long dash, short dash line. At an experimentally set potential

or threshold current the scan is reversed (scan 2) and is indicated by a solid line. After a time delay scan 3 (dashed line) is resumed from the

initial potential of scan 1 in the forward direction to a final potential.

The passive film stability is indicated by lower current dips during the

second scan. The rule of thumb suggests the plot be linear over several

decades for low current dips. However, if the second scan shows great

curvature, one can assume poor passive film stability. The alloy is said

to have less corrosion resistance the closer the third scan comes to retracing

the first scan.

b. Scratch Method

The scratch test technique, suggested by Pessall and Liu

[12], involves scratching the alloy surface at a constant potential and

measuring the current or current density as a function of time. This

technique is based on the concepts demonstrated in fig. 14 and suggests a critical potential, E^, that usually results in values between and

Ep. The figure illustrates how the authors expect the. potential dependences

of repassivation time and induction time for an alloy in a particular

environment to relate to a critical potential value, E^. Smooth surfaces

and rapid stirring would lead to greater induction times to break down the

protective film on the alloy. If this film is removed by scratching the

alloy surface below E^, the alloy would repassivate as shown in fig. 14.

33
The critical potential, E^, can be determined by the current

versus time curves at various potentials. When the potential is below E , the current observed upon scratching returns to the original values.

However, when the potential is above E^ there is little if any repassivation

and the current remains constant or increases. Figure 15 illustrates the

current transients above and below E^. A major advantage of the scratch

method is that it eliminates the induction time for localized corrosion to

be initiated. The results obtained by Pessall and Liu are dependent on the

alloy and the environment and can be characterized by a critical potential,

Eg, above which pits will nucleate and below which pits do not nucleate

(E = E). However, this value is usually less noble than E values obtained

s p

by potentiodynamic anodic polarization scans. In other words E^ separates

pitting tendencies from repassivation tendencies. Their results also

indicate that the protection potential is the same as the critical potential

Eg. Therefore, only one critical potential Eg, exists. Eg can also be

determined by gal vanostatic techniques; however, the results are difficult

to interpret in that large oscillations of the potential occur,

c. Tri bo-el 1 ipsometry

Tri bo-el 1 ipsometry is a technique developed to bypass the

initiation phase by abrading off the surface oxide and to measure the

repassivation of the exposed surface by ellipsometry to determine the film

growth kinetics [97,98]. The electrochemical cell with the retracting

abrasion wheel is shown in fig. 16. The initial phase of repassivation is

usually completed in less than 50 milliseconds so rapid withdrawal of the

abrasive wheel is desired. This equipment is designed so that el 1 ipsometric

measurement of film growth kinetics can begin within 10 milliseconds of

cessation of abrasion. Results of low carbon steel in nitrate, nitrite,

and buffered borate solution revealed significant oxide film growth had

occurred before the el 1 ipsometric measurements were taken. However, this

technique showed that in nitrite solution the rate of repassivation was greater and that after film rupture in the nitrate solution, the anodic

current went into metal dissolution. Therefore, tri bo-el 1 ipsometric measurements indicate film repair mechanisms as well as metal dissolution kinetics.

34
d. Pit Propagation Rate (PPR) Curves

Another technique to determine the rate of pit or crevice

propagation by avoiding long initiation times is the PPR curve technique

[99]. Figure 17 describes the PPR technique. It involves the following

steps:

(1)

Alloy sample is immersed in the solution for one hour

at E corr'

(2) Scan at 36 V/hr (10 mV/sec) to a preselected potential

between and (0.25 V in fig. 14) and hold this potential for 10 minutes.

Since this potential is below E , the current should give the rate of

P

general corrosion in the passive state.

(3) Continue the scan beyond E until a current density of

2 7 2 P

10 mA/cm (10 nA/cm ) is reached.

(4) Decrease the potential in a single step to the preselected potential between Ep and Ep^^^ (0.25 V in fig. 14) and hold this

potential for 10 minutes. The current density is a measure of the general

corrosion and the rate of pit growth.

(5) Decrease the potential in a single step to the original

free corrosion potential, to repassivate the pits. Hold this

potential for 5 minutes.

(6) Repeat (2) to ensure the current under passive conditions

have not changed significantly.

(7) Reduce the potential to E The general corrosion portion of the current is subtracted

from the total current found in (4) to obtain the shaded area in fig. 17

which is the average pitting current. The area due to pitting is determined

by microscopic examinations. Therefore, the pit propagation rate is the

average pitting current by the pitted area, and Syrett [100] refers to this

as the "average pit current density." Syrett' s results of 316L and 3 Mo

TRIP in arterial Tyrode's solution show similar pit propagation rate behavior

where as cyclic polarization scan results indicated significant differences

in that E^ - Ep^^^ might not be a good indication of crevice or pit propagation.

Also, cyclic polarization scans showed little difference of Ep^^^ for cold

worked and as cast 3 Mo TRIP but PPR curves demonstrated distinct differences

(fig. 18).

35
The scratch test and pit propagation rate curves avoid the

initiation time by removing the passive film or initiating pits at noble

potentials. Pessall and Liu [12] found light, small scratches to give

greater reproducibility. However, they assume the film will be broken in

the sites that are favored for pitting sites. The PPR curves scan to

potentials greater than the film breakdown potential to initiate pits, but

this value is dependent on the scan rate and surface effects of the alloy.

Nevertheless, both of these techniques along with rotating ring disk electrode

methods offer reasonably good quantitative information concerning propagation

rates. Initiation time can be estimated by studying passive film techniques

that shed light on the mechanisms of film breakdown.

e. Electrochemical Noise Measurements

This technique described by Bertocci and Kruger [37] examines

events leading to film breakdown, in particular the dynamic breakdown- repair

process preceding actual film breakdown. This process can be accomplished

by detecting and analyzing the fluctuations from the average values of the

current and electrode potential. Either voltage fluctuations under galvanostatic conditions or current fluctuations under potentiostatic conditions

can be monitored. The potentiostatic method was found to be preferred when

passive films are studied in that distinct energy conditions can be maintained

at the electrode/electrolyte interface. In studies of amorphous and crystalline

Fe-Ni-Cr alloy it was found that the amorphous compound resisted film

breakdown better because its homogenous film inhibited the dynamic breakdown

repair process to a greater extent. The electrochemical noise technique is

mainly useful for the detect! ofi of the onset of pitting and for the determination of Ep (see fig. 1).

f. Rotating Ring-Disk Electrode

Rotating disk and rotating ring disk electrodes can be used

to study corrosion phenomenon [101-104] and determine instantaneous corrosion

rates [106]. These corrosion rates were measured at open circuit as a function of rotation speed and temperatures. The RRDE collection experiment

involved locating the potentials of oxidation of the copper (I) product at

the ring and monitoring the partial anodic current in the dissolution of

copper at the disk. The ring current can detect directly surface alteration

without the averaging necessary in integral techniques such as various

potentiostatic methods. For further information about ROE and RRDE techniques

36
and results concerning corrosion mechanisms, see references 101 through 105.

Recent discussions with Ambrose [106] suggest a more appropriate technique

for determining propagation rates would involve a ring-ring electrode since

a ring disk electrode would probably have current transients depending on the location of the pits.

Standard testing procedures for cyclic potentiodynamic polarization

measurements for localized corrosion [107] and pitting or crevice corrosion

of metallic surgical implant materials [108] are available by the American

Standard Testing Measurements organization.

Significance of Current Understanding to Testing—Concluding Remarks

The current understanding of the mechanisms of pitting and crevice

corrosion we have described, incomplete and controversial as they are, never- theless provide valuable insights into test methods that could be useful in

assessing the life of metallic containers for nuclear waste. Out of our understanding of localized corrosion processes, four significant aspects can

be identified that appear to be of the utmost importance for the development

or choice of more reliable test methods:

1. The Existence of an Induction Period

The existence of an induction period, whatever the processes are that

control it, is by its very existence a measure of the time for pitting or crevice corrosion to initiate. Thus, a test that can measure the induction

period will also measure the time for initiation to take place. While such

tests exist (one simply holds a specimen above Ep by means of a potentiostat

and measures the time required to observe a rapid increase in anodic current),

induction times can be too long (greater than a year) to be measured in a reasonable test period. To our knowledge, there is no way one can extrapolate

test data to get a measure of the induction time and, hence, the time for

the initiation of localized corrosion. This implies that there is no way

to develop an accelerated test to measure, by extrapolation or otherwise,

the time required for localized corrosion initiation. This is, however,

not so for the propagation process. Once pitting or crevice corrosion is

initiated, test methods, such as the PPR method described earlier, can give

a value for the rate of propagation. This value, as imperfect as it may

be, can then be used to estimate the time required for penetration of the

contai ner wal 1

37
A question must now be posed, what does one do if the rate controlling

step is initiation and not propagation? The only course of action that we can suggest is the one that proposes the use of tests that eliminate the

initiation stage and measures the propagation rate only. If this propagation

rate is unacceptably high, critical potential measurements, which determine

conditions where initiation is not possible or, at least, probably occurs after an acceptably long induction period, must be used as a testing procedure. If one can be assured that the potential of the alloy used in the

environment in which the container will be emplaced is less than the critical

potential determined for that environment, our current understanding would

imply that the time for initiation would not be exceeded during the required

containment period.

2. The Existence of Critical Potentials

The importance of the existence of critical potentials, especially when

initiation processes control container life, has been discussed in item 1.

The major mechanistic and, therefore, testing problem associated with

critical potentials is that the literature is in a confused state providing

many "critical" potential values and procedures for their determination. In

defining many of these values, concepts of initiation and propagation have

been intertwined. For instance, the critical potential, E , in potentior dynamic scans is supposed to represent the value at which pits nucleate. Yet,

this value is dependent on the scan rate and therefore does not give a true

representation of initiation time for that or other potential values. The

protection potential is the potential below which repassivation occurs during

the experimental conditions. Therefore, a better understanding of the localized

corrosion process is needed to define more clearly the important critical potentials and how they may be obtained experimentally. With this in mind, it should

be possible to develop accelerated test methods that reflect real time values

for the initiation time and propagation rates for the localized corrosion of

alloys in a specific environment.

3. The Role of Repassivation

One key element in determining critical potentials is the presence or absence of a repassivation process and, quantitatively, the rate of such

a process. As pointed out earlier, some of the definitions of a critical

potential make E the potential below which repassivation takes place at

a high rate. Some theoretical models reserve for it an important role in

38
breakdown, i.e., initiation, processes. Propagation proceeds when repassivation can no longer choke off initiation. Test methods such as the scratch

test, triboel 1 ipsometry, and others, therefore, evaluate the possibility for

breakdown of passivity and propagation of localized corrosion of a given

alloy in a relevant environment (discussed in item 4) at a given potential

by measuring the rate of repassivation. If this rate is coupled to Ep, i.e., a high rate below Ep, a low rate above, its importance to testing for container

life is evident,

4. The Role of the Local Environment

Since most pitting and crevice corrosion mechanisms involve a change in

the local environment in a pit or crevice, it is reasonable to assume that

this environment should be the most relevant test environment. Such an assumption should, however, be established. The first task in establishing

the validity of the local environment as a test medium, actually an accelerated

test medium, is to characterize the local environments that develop in pits

and/or crevices for a given alloy in a given ambient environment. Once the

composition of these environments is known and the validity of their use as test media is established, it would seem reasonable to assume that such

environments would be the most relevant ones to use to measure critical

potentials, pit or crevice propagation rates, or repassivation rates.

In summary, a better understanding of the mechanisms of localized

corrosion can be used to suggest currently used tests or new tests to be

developed that can provide useful estimates of metallic nuclear waste

container life.

39
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40
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[49] B. E. Wilde, in ref. [2], p. 342.

[50] S. Smialowska and M. Czachor, in ref [2], p. 353.

[51] J. O'M. Bockris and A. K. N. Reddy, "Modern Electrochemistry," Vol.

2, Plenum Press, New York, 1970, p. 706.

[52] S. D. Argade and E. Gileadi, in "Electrosorption," E. Gileadi, Ed.,

Plenum Press, New York, 1967, P. 87.

[53] J. R. Galvele, S. M. de Micheli, I. L. Muller, S. B. de Wexler, and

I. L. Adams, in ref. [2], p. 580.

[54] H. S. Parker and C. A. Harding, J. Amer. Ceramic Soc., 583

(1970).

[55] J. Kruger and J. R. Ambrose, NBS Report, NBSIR 74-583, September

1974.

[56] K. J. Vetter and H. H. Strehblow, in ref. [2], p. 240.

[57] W. J. Lorenz and H. Fischer, Electrochemica. Acta. , V[, 1597 (1966).

[58] T. P. Hoar, J. Electrochem. Soc., 1J7, 17c (1970).

[59] C. L. McBee and J. Kruger, in ref. [6], p. 131.

[60] J. M. West, "Basic Corrosion and Oxidation," Ellis Horwood, Chichester,

1930, p. 114.

[61] T. Kodama and J. R. Ambrose, Corrosion, 33, 155 (1977).

42
[62] J. Kruger, J. Electrochem. Soc. , 106, 736 (1959).

[63] J. Kruger, J. Electrochem. Soc., 210, 654 (1963).

[64] M. Jani k-Czachor and Z. Szklarska-Smialowska, Corros. Sci., 8, 215

(1968).

[65] A. Szummer, Z. Szklarska-Smialowska, and M. Janik-Czachor, Corros.

Sci. , 8, 827 (1968).

[66] I. L. Rosenfeld, Korrzija i Zachita Metal lor, Lokalnyje Korrosionnyje,

Izd, Metallurgia (1970).

[67] A. Randak and F. W. Trautes, Werkstoff und Korrosion, 97 (1970).

[68] A. Szummer and M. Janik-Czachor, Brit. Corros. J. , 9, 215 (1974).

[69] Z. Szlarska-Smialowska, A. Szummer, and M. Janik-Czachor, Brit.

Corros. J. , 5, 159 (1970).

[70] A. P. Bond, G. F. Bolling and H. A. Domian, J. Electrochem. Soc.,

113 , 773 (1966).

[71] W. Schwenk, Werkstoff und Korrosion, 21> 791 (1968).

[72] T. P. Hoar, Corros. Sci., 7, 355 (1967).

[73] J. L. Rosenfeld and I. S. Danilov, Corros. Sci., 7, 129 (1967).

[74] J. R. Galvele, J. Electrochem. Soc. , 123 , 464 (1976).

[75] U. F. Franck, Werkstoff und Korrosion, 9, 504 (1958).

[76] K. J. Vetter and H. H. Strehblow, Ber. Bunsenges. physik. Chem. 74

1024 (1970).

[77] H. H. Strehblow, Keynote lecture, Toronto meeting. National Association

Corros. Engrs. , 1981.

[78] H. Kaesche, Z. physik Chem., N. F. , 34, 87 (1962).

[79] T. R. Beck, in ref. [2], p. 644.

[80] J. W. Oldfield and W. H. Sutton, Brit. Corros. J. , T3» ^3 (1978).

[81] I. L. Rosenfeld, in ref. [2], p. 373.

[82] G. Karlberg and G. Wranglen, Corros. Sci., Jl* 499 (1971).

[83] M. G. Fontana and N. D. Greene, "Corrosion Engineering," McGrawHill, New York, 1967, p. 41.

43
[84] D. B. Anderson, ASTM Spec. Techn. Publ. 1974, p. 231.

[85] F. P. IJsseling, Br. Corros. J. , (2), 51 (1980).

[86] A. Garner, Corrosion, 17 (3), 178 (1981).

[87] T. Suzuki and Y. Kitamura, Corrosion, ^ (1), 1 (1972).

[88] D. A. Jones and N. D. Greene, Corrosion, 25 (9), 367 (1969).

[89] Z. Szklarska - Smialowska and J. Mankowski, Corrosion Science, 18,

953 (1978).

[90] J. M. Defranoux, Corrosion Science, 3, 75 (1963).

[91] B. E. Wilde and E. Williams, J. Electrochem. Soc. , 118 , 1058 (1971).

[92] B. E. Wilde and E. Williams, Electrochim. Acta., 16, 1971 (1971).

[93] I. L. Rosenfeld and I. K. Marshakov, Corrosion, 115 (1964).

[94] H. P. Leckie, J. Electrochem. Soc., 1V7 (9), 1152 (1970).

[95] M. Jani k-Czachor, Werkstaffe Korros. 31^, 606 (1980).

[96] EG&G Princeton Applied Research Model 350A Corrosion Measurement

System.

[97] J. Kruger, "Electrochemical Techniques for Corrosion," R. Baboian,

Ed. , National Association of Corrosion Engineers, Houston, 1977, p. 35.

[98] J. R. Ambrose and J. Kruger, Corrosion, ^ (1), 30 (1972).

[99] B. C. Syrett, Corrosion, 33 (6), 221 (1977).

[100] B. C. Syrett, same as ref. 98 page 93.

[101] B. Miller, Acc. Chem. Res., 10, 54 (1977).

[102] Z. Zembura, Corros. Sci. 8, 703 (1968).

[103] Z. Zembura, Electrochim. Acta. , 1_0, 859 (1965).

[104] A. C. Riddiford, Adv. Electrochem. Electrochem. Eng. 4, 47 (1966).

[105] B. Miller and M. Bellavance, J. Electrochem. Soc., 119 (11), 1510

(1972).

[106] Private communications with John Ambrose.

44
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Figure Captions

Figure 1 Noise current spectra of 6061 A1 in borate buffer +

0.01 M NaCl. Potentiostatic conditions. Reference:

SCE. Averages over 64 spectra. (1) -700 mV, (2) -650 Me,

(3) -650 mV after 10 min, (4) -600 mV. From Bertocci and

Kruger [37].

Figure 2 The effect of potential on film growth in the absence and

-2

presence of chloride ion (10 M) in boric acid- sodium

tetraborate buffer solution (pH = 8.4). Potentials were

2 above (at -0.11 V). Film thickness was measured by

® 2

el 1 ipsometry. E^ is defined as Ep in the text. From

Ambrose and Kruger [18].

Figure 3 Stereographic triangle showing the effect of orientation of

iron grains on tendency of pit. From Kruger [62].

Figure 4 Schematic potential -pH diagram, a, pH of the bulk

solution; a-j a2> passive zone; pitting potential;

b, pH of the locally acidified zone. From Gravele [3].

Figure 5 Factors affecting crevice corrosion. From Oldfield and

Sutton [80].

Figure 6 Pit and crevice corrosion assemblies, (a) artificial

pit, ref. 87; (b) and (c) natural exposure tests, ref.

49; (d) Stern-Makrides assembly, ref. 88; (e) artificial

crevice assembly, ref. 87. Other designs can be found in

references 85, 86, and 89.

Figure 7 "Cyclic" potentiodynamic anodic polarization of 304

stainless steel in a ground water type environment at

26C. Scan rate 0.5 mV/sec (1.8 V/hr). Saturated calomel

electrode used as reference electrode.

46
Figure 8 "Cyclic" potentiodynamic anodic polarization curve of

three alloys containing a synthetic crevice at 25C,

aerated in 3.5 w/o NaCl solution, (a) Hastelloy C; (b)

Incoloy 825; (c) Carpenter 20Cb3. From Wilde [49].

Figure 9 "Cyclic" potentiodynamic anodic polarization curves for

Fe-30Cr-3Mo alloy in nitrogen saturated 1 M NaCl at 25°C.

Sweep speed 0.17 mV/sec. From Wilde [49].

Figure 10 Effect of potential scan on the value of Ep for 304

stainless steel in 0.1 M NaCl. From Leckie [107].

Figure 11 "Cyclic" polarization behavior of 430 stainless steel in

1 M NaCl demonstrating the striking effect of pit propagation

on E^. From Wilde [49].

Figure 12 A schematic representation of (a) a "cyclic" potentiodynamic

anodic polarization curve; (b) a gal vanostatic curve; (c)

a potentiostatic curve with a passivated surface; (d) a potentiostatic curve with an activated surface. From

S. Smialowska and M. Czachor [50].

Figure 13 A corrosion behavior diagram plot of 304 stainless steel,

forward scan (---), reverse scan (solid line), third

scan after a time delay (dashed line). From EG&G Princeton

Applied Research [97].

Figure 14 Potential versus time plot of scratch test illustrating the

critical potential, E^, as it relates to the repassivation

time and the induction time. From S. Smialowska and

M. Czachor [50].

Figure 15 Scratch test current-time curves for specimen held potentiostatically above Ep (pitting or breakdown) and below Ep

(repassivation). From J. Kruger [98].

47
Figure 16 Electrochemical cell used in tri bo-el 1 ipsometry studies

of repassivation kinetics. From J. R. Ambrose and J. Kruger

[99].

Figure 17 Schematic of potential-time and current-time cycles used

to determine the pit propagation rate curves. From Syrett

[101].

Figure 18 Effect of cold work on PPR curves for 3Mo TRIP. From

Syrett [10].

48
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NBS»114A (REV, 2-BCI U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

1. PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No, 3. Publication Date

SHEET (See instructions) NRSIR 82-2477 April 1982

4. TITLE AND SUBTITLE

Current Understanding of Pitting and Crevice Corrosion and Its Application to

Test Methods for Determining the Susceptibility to Such Corrosion of Nuclear

Waste Metallic Containers •

5. AUTHOR(S) ; Jerome Kruger and V. Kay Hardman

6. PERFORMING ORGANIZATION (If joint or other than NBS, see Instructions)

national bureau of standards

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

7. Contract/Grant No.

8. Type of Report & Period Covered

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street, City. State. ZIP)

Nuclear Waste Management Program

Department of Energy

Germantown, Maryland

10. SUPPLEMENTARY NOTES

Document describes a computer program; SF-185, FIPS Software Summary, Is attached.

11. ABSTRACT (A 200-word or less factual summary of most significant Information. If document Includes a significant

bibliography or literature survey, mention it here)

Localized corrosion is the most damaging and prevalent type of corrosion that

can affect metallic containers used in nuclear waste repository sites. Mechanisms

of pittin and crevice corrosion and the techniques used in elucidating these

mechanisms are discussed in this review which includes a literature survey.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons)

Accelerated testing; crevice corrosion; electrochemical techniques, localized

corrosion, localized corrosion mechanism; pitting

13. AVAILABILITY

y ty I Unlimited

I I For Official Distribution. Do Not Release to NTIS

Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C.

20402.

[y~| Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

69

15. Price

$9.00

USCOMM-DC 6043-PBO
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APRIL2013

Current understanding of pitting and crevice corrosi(

Kruger, Jerome;Hardman, V. Kay

QC100.U56 no.82-2477 1982

NIST Research Library

[174] cuiTentunderstan8224krug

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