NBSIR 76-1170

The Role of Passive Film

Growth Kinetics and Properties

in Stress Corrosion and Crevice

Corrosion Susceptibility

J. Kruger and J. R. Ambrose

Corrosion and Electrodeposition Section

Metallurgy Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

November 1976

Technical Summary Report Number 7

Prepared for Office of Naval Research

Department of the Navy

Arlington, Virginia 22217

NBSIR 76-1170

THE ROLE OF PASSIVE FILM

GROWTH KINETICS AND PROPERTIES

IN STRESS CORROSION AND CREVICE

CORROSION SUSCEPTIBILITY

J. Kruger and J. R. Ambrose

Corrosion and Electrodeposition Section

Metallurgy Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

November 1976

Technical Summary Report Number 7

DISTRIBUTION OF THIS DOCUMENT IS UNLIMITED

Prepared for Office of Naval Research

Department of the Navy

Arlington, Virginia 22217

U.S. DEPARTMENT OF COMMERCE. Elliot L. Richardson. Secretary

Edward O. Vetter. Under Secretary

Dr. Betsy Ancker-Johnson, Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS. Ernest Ambler. Acting Director

PART I

[To be included in the Proceedings of Symposium on Inhibitors for Localized

Corrosion, NACE National Meeting, San Francisco, March 1977.]

THE ROLE OF MOLYBDENUM AS AN INHIBITOR OF LOCALIZED CORROSION

ON IRON IN CHLORIDE SOLUTIONS

J. R. Ambrose

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

ABSTRACT

A series of Fe-xMo alloys, where x ranged from 0 to 6 weight

percent, were used to determine the effect of molybdenum concentration

on the repassivation kinetics of iron within a crevice. Results show

that concentrations of molybdenum greater than 5 weight percent are effective in increasing the repassivation rate. Resistance of these

alloys to crevice corrosion attack appears to result from the role of

molybdenum in the formation of a protective salt film following breakdown of passivity within a crevice.
Introduction

The term "Inhibitor for Localized Corrosion" has come to be asso- ciated with some additive which, when introduced into an environment,

improves the resistance of a particular metal to pitting, crevice, or

stress corrosion cracking. However, recent studies have shown that

certain alloying elements, molybdenum as a case in point, behave in a fashion similar to what has come to be expected for inhibitors added

directly to the solution. In other words, the mechanism by which the

localized corrosion processes is inhibited is the same whether the

element is alloyed Into the metal, or whether its anodic dissolution

product is added to the solution. Reports of such behavior have been

published for molybdenum, chromium, vanadium, silicon and several other

less common alloying elements (1-7). As a caterpillar can protect

itself when it generates its cacoon, more efficient protection of an alloy may be likely when the alloy generates its own inhibitors.

A study by Kodama and Ambrose (1) showed that the molybdate ion,

which is the anodic dissolution product of molybdenum, would inhibit

propagation of pits but not their initiation. The explanation of this

effect was based on the observation that these repassivated pits contained

measurable amounts of the element molybdenum possibly in the form of

FeMoO^. What was not readily available from these experiments was a detailed knowledge of how these pits repassivated; if a salt film was involved, how fast did it grow? To meet the need to examine the changes

occurring at the metal surface, a technique was developed (8) for

ell ipsometric measurements of optical changes inside the crevice created

by placing a quartz plate adjacent to a reflective metal surface. Using

this technique, the effect of introducing a crevice on film growth and

dissolution kinetics on iron could be measured and related to electrochemical changes.

This paper is concerned mainly with the role molybdenum plays in

iron repassivation kinetics and hence how repassivation rates could be

expected to control susceptibility to localized corrosion processes. It

is becoming increasingly recognized that repassivation is one of the

major basic processes affecting localized corrosion (9). To measure

repassivation kinetics in a crevice presents certain difficulties. For

example, cathodic reduction of surface film on iron within a crevice is

impossible in the presence of a reducible species such as Mo04

=

. Addition

of the ion just before repassivation is measured is also a difficult

procedure. Therefore, a series of binary Fe-Mo alloys was prepared with

the intention of comparing results from this study with those obtained

when molybdate was added directly to the solution. The work described

here involves repassivation of iron-molybdenum allo'ys in sodium borateboric acid solutions to which sodium chloride solution was added. These

studies, in addition to providing some answers to the question of role

of molybdenum as an inhibitor, also point to future work necessary to

evaluate inhibitor performance.

- 2 -
Experimental

Binary iron-molybdenum alloys of composition Fe-xMo, where x ranged

from 0 to 6 weight percent, were prepared from high purity materials

(see Tables I and II). The iron had previously been annealed in hydrogen

to remove reducible species. The alloys were homogenized for 72 hours

at 1000 °C and under a vacuum of not less than 5 x 10"7 torr, then

furnace cooled. Cylindrical specimens were machined from the castings

using an electrical discharge spark erosion process and then mounted in an epoxy resin. The mounted specimens were then drilled and tapped for

electrode assembly. A short section of glass tubing was cemented with

epoxy resin into a hole drilled completely through the electrode and was connected via a flexible piece of surgical tubing to the salt bridge of

the reference electrode assembly. This particular arrangement permits

measurement of electrode potentials within a crevice (Fig. 1). The

remainder of the crevice electrode assembly (Fig. 2) was the same as that used in previous studies ( 10 ).

Specimens were polished using silicon carbide metallographic

papers, then polished on a wheel using 6 ym and finally 1 ym diamond

pastes. Following polishing, specimens were ul trasonical ly cleaned with

ethyl alcohol, rinsed with spectro-grade methanol, dried in air, then

mounted in the teflon retaining ring. The entire crevice electrode

assembly, with the quartz disk in place (Fig. 3) was stored in a vacuum desiccator until ready for exposure to the solution.

All experiments were performed using solutions prepared from

reagent grade chemicals and distilled water (7 x 10“' ohm

-

! cm~l). A

buffer solution containing 9.3 g H

3 BO 3 and 14.3 g ^B^ • 10 H

2 O per

liter (pH 8.4) was used as a stock solution. Solutions containing

chloride ion were prepared by adding an aliquot of 1.0 N NaCl solution

to the stock solution. The chloride ion concentration used in these

studies was 5 x 10-3 N in all cases.

Just prior to positioning the crevice electrode assembly in place

on the el 1 ipsometer, the crevice thickness was set by spacing the quartz

disk 0.25 mm from the metal surface using a teflon shim. After connection

to a potentiostat and transient detection equipment, the crevice electrode

was flushed with fresh solution through the salt bridge. Surface airformed films were reduced electrochemical ly at -758 mV SHE with the

potentiostat. After a period of approximately ten minutes, when the

cathodic current level had stabilized, ellipsometer readings and calibrations were made. The potential was then stepped to +242 mV SHE;

current and charge transients were recorded on an oscilloscope with

extended time transients being measured on a strip chart recorder. A

similar procedure was followed in those experiments where no crevice was present.

- 3-
Table I. Iron Composition

Impurity Content Impurity Content

Heat No. Element (weight percent) Element (weight percent)

C80022-1 c

- .006 Co .003

Mn ' .001 Cu .001

P .003 A1 .01

S .007 Sn .001

Si .008 Pb .0006

Ni .03 02 .026

Cr .01 h

2 .00006

V .004 n 2 .0004

W .01

Mo .001

- 4

-
Table II. Molybdenum Composition

Impuri ty

Element

Content

(wt.ppm)

Impuri ty

Element

Content

(wt.ppm)

Impuri ty

Element

Content

(wt.ppm)

0 < 10.0 Ca 0.2 Nb 3.0

C 40.0 Sc < 0.4 Ag < 2.0

F < 6.0 V 0.5 Cd <20.0

Mg < 2.0 Cr 0.02 In < 1.0

A1 0.1 Mn 2.0 Sn < 1.0

Si 2.0 Fe 0.1 Te <15.0

P < 0.1 Co 0.04 Ta < 6.0

S < 0.6 Ni < 10.0 W 60.0

Cl 0.4 Cu 0.04 Au < 6.0

K 0.2 Zr < 10.0 Pb < 0.4

Others <0. Ippm

- 5

-
In each case, simultaneous measurement of decrease in ellipsometer

light intensity (proportional to film thickness) was made using an offnull technique (11). After each experiment, the electrode was disassembled, the specimen repolished, solutions changed, and the entire

process repeated. Results are based upon an average of two consecutive

runs.

Resul ts

_

"3 In 5 x 10 N buffered chloride solutions, repassivation at 242 mV

SHE following cathodic film reduction at -758 mV SHE took place in two

distinct stages. During Stage I the anodic current generated as a result of the potential step was found to decrease at a rate proportional

to the concentration of molybdenum in the alloy when the alloy was exposed

within a crevice. Current decayed until a time interval proportional to

the molybdenum concentration in the alloy had elapsed. At this point in

time, referred to as t ss , the current ceased to decay, and what will be

called Stage II began. During Stage II, metal dissolved at a fixed rate

inversely proportional to the concentration of molybdenum in the alloy.

In Fig. 4 the anodic charge consumed during the repassivation for

three of the materials used in this study is plotted as a function of

the square root of time; no crevice was present. During Stage I, where

the data points fall along straight lines, current is decreasing.

Steady state current levels are indicated by deviations of data points

from the straight lines. The time at which these deviations appear is

labeled t §s in the figure. The effect of increased molybdenum content

in the alloy is evident from the observation that the decrease in the

rate of consumption of charge goes from 10“1 coulombs sec“V2 for Fe-1

Mo to 7 x 10"2 coulombs sec

- l/2 for Fe-3 Mo to 3.1 x 10^2 coulombs sec" ' for Fe-5 Mo. Moreover, t §s increases from approximately 200 milliseconds to over 1000 milliseconds as the molybdenum concentration in the

alloy is increased from 1 to 5 weight percent. Finally, the Stage II constant current level decreases from 365 to 7.3 milliamperes as the

molybdenum content increases from 1 to 5%. Interestingly enough, for

the same Stages I and II, values of current intermediate between those

for Fe-3 Mo and Fe-5 Mo were found for pure iron (5.3 x 10

- 2 coulomb

sec

- and 11.8 milliamperes, respectively). These results are tabulated

in Table III.

For alloys exposed to the crevice, somewhat more irregular behavior

was measured. Within a crevice, no repassivation of pure iron occurred

since the current did not decay from its initial peak value. The effect

of molybdenum addition is shown in Table IV. Although pronounced decreases

in current level are evident as the Mo concentration is increased to 5 weight percent, an anomolous behavior is observed for the Fe-3 Mo alloy.

Here, current levels are somewhat higher than the values measured for

the Fe-1 Mo alloy.

- 6 -
Table III. Non-Crevice Repassivation Kinetics

% Mo

Stage I Dissolution Rate (xl(r),

- 1/2 coulomb-sec

Stage II Dissolution Rate

mA

0 53 11.8

1 100 36.5

3 70 24.6

5 31 7.3

Table IV. Crevice Repassivation Kinetics

Stage I Dissolution Rate (xlO^) Stage II Dissolution Rate,

% Mo

-1/2 coulomb-sec - mA

0 83

1 115 7.6

3 240 15.1

5 72 6.3

* no current decay from peak value observed.

One consistent observation could be made, however. Introduction of the

crevice delayed the onset of Stage II in each alloy studied. In other

words, the current would continue to decrease for longer time intervals

when the crevice was in place than when it was not present. This leads,

of course, to lower steady state current values (Stage II) for a given

alloy composition inside a crevice as compared to the tss found with no crevice. Finally, as shown in Fig. 5, current levels, hence rate of

charge consumption for the Fe-5 Mo alloy during Stage I, were higher

inside the crevice than outside. The phenomonology was the same for the

other alloys studied.

Ellipsometer measurements were made simultaneously with the electrochemical ones. The off-null technique was used in which optical changes

which occur at the metal surface give rise to changes in the photocurrent

intensity measured during the repassivation transient. Results for

three alloys, when not exposed to a crevice, are shown in Fig. 6. Here,

changes in intensity, which are proportional to AP (change in polarizer

angle) and also to film thickness, are plotted against the log of the

square root of time. Although rates of change of surface optical

properties are about the same, the relative magnitude of aP is less for

the higher molybdenum-containing alloy.

- 7-
Similar measurements for the material exposed within a crevice were

not as successful. Due to the electrode configuration in this particular

cell, the position of the counter electrode with respect to the specimen

resulted in most of the cell current being consumed at an area on the

working electrode surface which was closest to the auxiliary electrode.

This particular area happened to be out of view by the ell ipsometer. A new electrode design, which incorporates the counter electrode into the

specimen as a ring around the working electrode (Fig. 7), promises to

allow better current distribution to the working electrode and more

meaningful ell ipsometer measurements.

Pi scussion

1 . Role of mass transport in repassivation kinetics

The study by Kodama and Ambrose (1), in which additions of MoC^

-

to chloride containing solutions appeared to affect the propagation

stage of pitting, showed the effect of increased concentration of the

molybdate ion (Fig. 8). Repassivation of these pits was attributed to

the formation of a stable salt film which restricted metal dissolution.

Similar results, which also may be explained by the formation of a molybdate containing salt film, were obtained in this study.

Stage II repassivation kinetics as described here correspond to the

plateau areas in Fig. 8 where pits already initiated are beginning to cease to propagate. By using a crevice, however, it is possible to

magnify the area of localized corrosion to such an extent that more

careful examination of the influence of molybdenum (or its anodic

dissolution product, M0Q4

-

) is possible. By comparing what happens when

a crevice is present to what happens when it is absent makes it evident

that, for molybdenum to affect the repassivation kinetics of the material to which it is alloyed, it must be introduced into the environment

under conditions of severe mass transport limitations. We can see, for

instance, that when repassivation with and without a crevice are compared,

although the metal dissolution rate for unalloyed iron is some 7 times

higher during Stage II crevice repassivation (Tables III and IV), the

effect of alloyed molybdenum1 is to decrease the dissolution rate during

Stage II over that measured without a crevice. This increase in re- passivation efficacy occurred even when current levels were significantly

higher during Stage I repassivation (Fig. 5). What this suggests is

that unless the anolyte becomes very concentrated with the molybdate

ion, the dissolution rate of any salt film containing this ion will be

of such magnitude to preclude its stable existence for long enough

periods of time to afford protection. This suggestion also explains

why, in the previous study (1), molybdenum was only detected within

those pits which had initiated, but which subsequently had failed to

propagate. Restricted geometries of these pits provided the limited mass transport conditions necessary to build up a high concentration of

molybdate ion, to allow formation of a stable salt film, and then to

prevent dilution of the molybdate solution in these newly passivated

pits by diffusion of the ion into the bulk solution. Consistent with

- 8-
the idea of formation of a salt film is the experimental evidence that

the degree of repassivation is related to the concentration of the

molybdate ion in solution during the pitting of iron (1) and to the

magnitude of anodic dissolution found during crevice corrosion in this

study. Several critical experiments remain to be done to test this

suggestion whether an insoluble salt film may be forming. Accurate

anolyte solution analysis with respect to iron and molybdate ions is

needed to establish concentration profiles within the crevice volume.

Secondly, surface profile analysis of the exposed metal must show that

molybdenum can indeed be enriched under those conditions of limited

solution mass transport. The results of these particular experiments

would go a long way to explain the results of researchers (12) who,

using Auger Electron Spectroscopy techniques, fail to observe surface

molybdenum enrichment following anodic dissolution in ferritic stainless

steels.

2. What affects salt film formation ?

Accepting for the moment the idea that formation of salt films of

the sort suggested by Kodama and Ambrose is possible, several questions

remain to be answered:

a. What parameters determine the effectiveness of a particular

material in enhancing repassivation behavior? Are they related

to initial rate of formation, overall uniform thickness, stability

under a wide range of solution pH, potential, chloride ion

concentration?

b. What is so special about molybdenum? Can the properties of the

alloyed metal and its dissolution products be used to predict

performance of other suggested inhibitors - whether they be in the

form of an alloy or as an additive to the environment?

Some answers come from examination of data obtained during Stage I repassivation. If anodic current decay can be related to thickening of

a protective surface film, then it may be supposed that so long as the

film grows, current will continue to decrease. From Table III it can be seen that the rate of current decay is related to % Mo in the alloy. In

Fig. 4 it is shown that the time during which current continues to decay

is also proportional to the molybdenum concentration in the alloy.

Thus, for the case where no crevice is present, it seems clear that

molybdenum can contribute to improved repassivation behavior by stimulating formation of a thicker, more protective surface film. However, we should temper our conclusions by examining what the ellipsometer

measurements show with regard to what is happening on the metal surface.

From Fig. 6, it can be seen that these measurements do not show this to

be true - in fact the inverse relationship is demonstrated. Such

results are not surprising, however, when one remembers that kAP is

proportional not only to film thickness but also to degree of surface

roughening (pitting) and to accumulations of non-protecti ve corrosion

- 9-
product deposits. What Fig. 6 does show then, is that deterioration of

the surface is restricted as the molybdenum concentration is increased.

One other point should be kept in mind. The requirement of solution

mass transport limitation already discussed is met only within the pits

initiated on the surface, or in the crevice used in this study. No

ellipsometer currently available can measure the optical changes that occur within these pits.

Current densities measured during the course of this study were

found to range from 1-100 mA/cm3 . When we distribute this current in

proportion to the composition of the alloy, one can get an idea of the

rate of build-up of molybdate ion as approximately 10" 3 equivalents

liter"! sec"!. After about 1 second, the concentration of molybdate in

this particular crevice volume would be 10" 3 N, or roughly equal to the

minimum concentration of MqOa" necessary to inhibit pit propagation on iron in chloride solutions of 10" 3 N (1). It should be remembered,

however, that sufficient molybdate must be available to produce a salt

film of critical thickness before the solution has deteriorated too far.

Once pH and chloride ion concentrations have reached levels where the

salt film remains unstable, propagation of pits or crevice attack

cannot be stopped. Hence, the fact that low molybdenum composition

alloys do not improve corrosion resistance of pure iron, no matter how

long one waits to build up the concentration of M0 O4

" ion in the re- stricted volume of solution.

3. Is molybdenum always beneficial ?

A word of caution is necessary before completely accepting the

existence of effects of molybdenum in iron. In Table III it can be seen that the addition of ]% molybdenum to iron actually results in a decrease

in the Stage I repassivation rate as well as in increase in the steady

state corrosion during Stage II. In fact, no real improvement over the

performance of pure iron is noticed until the molybdenum concentration

in the alloy reaches 5 %. This effect is magnified when a crevice is

present (see Table IV) since we are obviously dealing with a much larger

volume of solution in which mass transport is limited. This observation

emphasizes the concerns of question 1. It is intuitive that additions

of molybdenum should change the properties of a surface film. But which

way? It stands to reason that if a film on iron is protective enough

under a given set of conditions, then it is useless to try to affect the

properties of that film so long as those conditions prevail. But when

conditions are altered such that the film becomes non-protective, as it

is during pitting or crevice corrosion attack, it is then that an improvement in the protection properties of a surface film is required.

Thus, during early stages of repassivation of iron when solution pH and

composition are such that if insufficient concentrations of molybdate

are available, the film formed will not be as protective as if no molybdenum were present at all. It also appears, all other things being

equal, that a critical concentration of molybdenum, whether present in

the metal or in solution, is necessary to provide a more or less uniform

- 10 -
film of a critical thickness. It appears that the minimum molybdenum

concentration for these Fe-Mo binary alloys is around b%. Introduction

of other alloy constituents would probably have a pronounced effect on the value (10). This study showed that increased concentrations of

chromium improved upon the effect of addition of molybdenum.

4. What remains to be done ?

It is hoped that verification of the presence of the salt film and

its thickness dependence on molybdenum concentration in the alloy during

crevice repassivation will come from use of the new specimen design

(Fig. 7). These results coupled with solution microanalyses, and

surface profile analysis with ESCA or Auger spectroscopy will go a long

way towards answering a great many of the questions which have arisen

from an interpretation of the results obtained in this work.

Conclusions

1. Additions of less than 5 l molybdenum are detrimental to the

repassivation kinetics of iron with a crevice present or not.

2. Repassivation of these Fe-Mo binary alloys takes place in two

stages - Stage I, where current decreases are measured, and Stage II,

where a steady state current level is observed. Stage I corresponds to

salt film growth. Stage II to constant salt film thickness.

3. Introduction of a crevice results in an increase in the metal

dissolution rates for a given alloy composition during Stage I repassivation.

However, Stage II steady state corrosion rates are decreased within a crevice as the molybdenum concentration in the alloy is increased.

4. Results are interpreted on the basis of formation of a protective

film during repassivation. A critical thickness of film formed within a critical time interval will determine whether the particular localized

corrosion process will propagate or not.

5. Solution composition analysis, surface profile analysis and

improved ellipsometer measurement techniques should resolve many of the

interpretations presented as a result of repassivation kinetics measurements of these iron-molybdenum binary alloys.

Acknowledgment

I am grateful to the Office of Naval Research, who supported this

work under Contract No. NANR 18-69, NR 036-082.

- 11 -
Fig. 1 . Fe-3Mo alloy mounted in epoxy resin. The reference electrode

salt bridge can be attached to the glass tubing which

protrudes from the rear of the electrode.

- 12 -
Fig. 2. Components of the crevice electrode. Shown from top right

to bottom left are the working electrode attached to a plexiglass support, a stainless steel rod for electrical

connection to the specimen, the quartz disk and teflon

adapter used for making the crevice.

- 13 -
- 14 -
0

Dissolution Kinetics no crevice

5 x IO' 3 N Cl"

0

1/2 Fig. 4. Plot of charge consumed, Q, versus square root of time, ms for Fe-xMo (x=l,2,3,) during repassivation at 242 mV SHE in

the sodium borate-boric acid buffer; chloride ion concentration

= 5 x 10"3 N.

- 15 -
t l/2

, ms l/2

Fig. 5. Comparison of dissolution kinetics of Fe-5Mo within and without a crevice. Plotted are changes in consumed charge, Q, versus square root of time, msV2, Repassivation occurs at

242 mV SHE in the sodium borate-boric acid buffer; chloride

ion concentration = 5 x 10~3 N.

- 16 -
10 100

'/2 log,0

t (ms)

Fig. 6. Plot of ellipsometer variation, kAP versus logarithm square

root of time, 1/2 logiQt, in ms, for Fe-xMo (x = 1,2,3,) with

no crevice in place. Repassivation occurs at 242 mV SHE in

the sodium borate-boric acid buffer; chloride ion concentration

= 5 x 10-3 N.

- 17 -
Fig. 7, Fe-3Mo alloy with encircling platinum electrode mounted in

epoxy resin.

- 18 -
CURRENT,

mA

TIME, s

Fig. 8. Plot of anodic current decay against the logarithm of time in

seconds for 4 concentrations of sodium molybdate, 0, 5 x 10~3,

10"2, 3 x 10~2 M. Repassivation occurs at 442 mV SHE in the

sodium borate-boric acid buffer, chloride ion concentration = 10~2 m [from (1 )]

- 19 -
REFERENCES

1. T. Kodama and J. K. Ambrose, submitted to Corrosion, 1976.

2. N. D. Tomashov, G. P. Orernova, and 0. N. Marcova, International

Conference on Localized Corrosion, Williamsburg, Va., 1971 NACE-3

(Houston, 1974) 363. i i

3. W. Robertson, J. Electrochem. Soc. 98, 94 (1951).

4. M. Pryor and M. Cohen, J. Electrochem. Soc., 100, 203 (1953).

5. H. Goldschmiedt, Paint Manuf., 45, 22 (1975).

6. Y. Fukase, K. Osozawa, R. Aemoto, Japan - U. S. Seminar on the

Physical Metallurgy of Heat Resisting Alloys, Diamond Park, New York

(1972).

7. G. A. Lizlovs, Corrosion, 32, 263 (1976).

8. J. R. Ambrose and J. Kruger, Corrosion, 28, 30 (1972).

9. J. C. Scully, The Theory of Stress Corrosion Cracking in Alloys,

NATO Science Committee, Brussels, (1971), 1.

10. J. R. Ambrose and J. Kruger, Technical Summary Report No. 5, NBSIR

74-583 (1974), 1.

11. J. J. Carroll and A. J. Melmed, Surface Science, 16., 251 (1969).

12. J. B. Lumsden and R. W. Staehle, Scripta Metal lurgica, 6, 1205

(1972).

- 20-
PART II

REPASSIVATION KINETICS OF MECHANICALLY DEPASSIVATED

METAL WITHIN A CREVICE

Introduction

One of the problems associated with any study into what factors are important in controlling susceptibility to crevice corrosion is that one must wait out some incubation period before measurable attack has begun.

And it can well be that during this period subtle changes are occurring

within the crevice - the solution, the films, the metal itself - that

play a critical role in what will happen following some breakdown in

passivity. For this reason, a technique which allows metal exposed to a crevice to be quickly depassivated after any elapsed exposure interval

is necessary if one wishes to know what is really happening during the

initial stages of crevice corrosion attack.

i This work deals with the development of a technique which utilizes

a pointed sapphire rod to scratch the surface of a metal electrode

exposed within a crevice. Following this mechanical depassivation,

repassivation kinetics can be measured using conventional electrochemical

techniques. To date, first generation alterations in the original

concept of the device have been made; these changes have led to a new type of crevice electrode. Other results have led to criteria of

performance which will be carried over into any second generation

design.

Experimental

The device used to measure repassivation kinetics of metal within

a crevice basically consists of a female standard taper ground glass

joint, the side arms of which have been modified for use as the elements

of an electrochemical cell. In these side arms are contained a reference electrode bridge, an auxiliary platinum electrode, and solution

addition/drain stopcock. In the same arm as the counter electrode is

the sapphire tip assembly. Pressure to push the tip against the specimen

is applied with a stainless steel spring, held in place by a Swagelok

cap. The assembled cell is shown in Fig. 1.

The original crevice electrode is shown in Fig. 2. It is a 304

stainless steel plug machined to the specifications of a standard 19/38

taper and designed to fit the glass cell. This electrode as shown has

been masked off with glyptal varnish to reduce the area of electrode

actually in contact with the crevice electrolyte.

A new electrode design has been used in the studies presently

underway. It incorporates a metal electrode section sandwiched between

two pieces of teflon. The unassembled and assembled electrodes are shown in Figs. 3 and 4.

- 21 -
The dimensions of the crevice are controlled by the vertical

position of the electrode in the crevice. This position is adjusted

using the wheel around the threaded central shaft of the electrode.

Mechanical depassivation is accomplished by rotating the electrode

assembly in the crevice. Using a potentiostat, current decay during

repassivation can be recorded on an oscilloscope; longer transients are more conveniently measured on a strip chart recorder.

Repassivation kinetics measurements of AISI 304 stainless steel in

sodium chloride solutions are currently in progress.

- 22 -
Fig. 1 . Assembled cell for determination of repassivation kinetics for

mechanically depassivated material within a crevice.

- 23 -
Fig. 2. Original machined 304 stainless steel electrode. Most of the

exposed surface has been coated with laquer to reduce contact

with the environment.

- 24 -
I

Fig. 3. Unassembled electrode used in crevice repassivation studies.

Shown are sapphire scratching device, metal specimen and the

teflon sections used to assemble the electrode.

- 25 -
Fig. 4. Assembled crevice electrode used in crevice repassivation

measurements.

- 26 -
PART III

REPASSIVATION KINETICS OF SIMULATED CHROMIUM DEPLETED

ZONES IN AUSTENITIC STAINLESS STEELS

Introduction

Susceptibility of austenitic stainless steels to intergranular

stress corrosion cracking (ISCC) during exposure to high temperature,

high purity nuclear reactor water systems has been attributed to the

presence of heat affected areas where the steel has been sensitized

during welding (1). Sensitization has been said to result when formation of chromium carbide deposits (C^gCg) in the grain boundaries leave

an adjacent area of rather active zones where chromium has been depleted

(2)

. Although extent of sensitization has been found to be related to

susceptibility of this alloy to ISCC for materials which have been

furnace sensitized, there has been great difficulty in establishing

similar definite relationships for material which has been sensitized

through welding alone (3). In other words, sensitization is a necessary

but not sufficient condition to account for the susceptibility of

austenitic stainless steels for ISCC. If, then, conventional time to

failure stress corrosion cracking tests are not reliable indicators of

material susceptibility, a need exists for a technique by which such

susceptibility can be ascertained.

It has been suggested that a relationship exists between the rate

at which a bare metal surface will repassivate and the susceptibility of

that metal to stress corrosion cracking in that particular environment.

In a previous study we have demonstrated the existence of that relationship

for AISI 304 stainless steel as a function of pH and applied potential

(3)

. Since chromium depletion is thought to result in a decrease in the

protectiveness of the passive film on this particular steel, it was thought that the technique used then would be equally applicable now. This technique, tribo-ell ipsometry, permits simultaneous measurement of

film growth and current decay kinetics during the repassivation transient

which follows surface film removal by mechanical abrasion (4).

This study will focus on the effect of systematic chromium depletion

on repassivation kinetics of a series of Fe-lONi-xCr where x varies from

18 to 3 weight percent. Along with these measurements, cathodic reduction

kinetics studies will be made on a specimen of Cr^oCg, the precipitated

carbide found in grain boundaries of sensitized stainless steels.

To date, only the specimen preparation stage of the study has been

completed.

Experimental Simulated Chromium Depleted Zones

To determine what effect selective depletion would have on the rate

at which these materials would repassivate, a series of ternary Fe-10

- 27 -
Ni-xCr alloys were prepared in which the chromium concentration was varied betwen 3 and 18 weight percent. These prepared alloys are intended to simulate the chromium depleted regions produced during the

sensitization process. Considerable effort was made to retain the

austenite structures in all alloy compositions through variation in

cooling rates and in carbon and nickel compositions, but the structure

of the final alloy series consists 1 of a martensite-like phase in a ferritic stainless steel matrix. ' 1

The compositions prepared were as follows:

Fe 18 W/0 Cr 10 W/0 Ni

Fe 15 W/0 Cr TO W/0 Ni

Fe 12 W/0 Cr 10 W/0 Ni

Fe 9 W/0 Cr 10 W/0 Ni

Fe 6 W/0 Cr 10 W/0 Ni

Fe 3 W/0 Cr 10 W/0 Ni

only specimen austenitic in structure

Stoichiometric compositions of the above alloys were carefully

weighed out using the highest purity materials available.

About sixty grams of each alloy was weighed out and then melted in

the arc furnace under titanium gettered argon. The alloy was melted a minimum of four times on each side with the sides being alternated

consecutively. After a total of eight meltings the alloy is in the

shape of a hemispherical button and is then arc cast into an approximately

square rod.

After the alloy was cast into a rod, it was removed from the arc furnace and mechanically deformed and shaped into a round rod by swaging

to the desired diameter.

The alloy was then heat treated for the purpose of homogenization,

stress relief and recrystal 1 ization of the as cast structure.

The heat treatment went as follows. The specimens were sealed in an evacuated vycar tube heat treated at 1000 °C overnight and quenched

in ice water in the morning.

Ambient temperature repassivation kinetics measurements using a modification of the tribo-el 1 ipsometric technique are presently in

progess.

Chromium Carbides, Cr^C^

This material was prepared by F. S. Biancaniello of the Alloy

Preparation Laboratory, Metallurgy Division, National Bureau of Standards.

- 28 -
I

!

Cr 23C 6 has a cubic > d84 type °f structure, a=10-638kX, Dm

=6‘97

1 g/cm3 M=4, Fm3m. This phase was earl ier' referred to as Cr4

C.

Stoichrometric quantities of chrome pieces (varying from -100 mesh

j up to about 0.5 grams) and carbon (lamp black fine powder) were weighed

out and placed in an alumina crucible which was placed in the Balzers

vacuum, melting and casting induction furnace for overnight pumpdown.

The following day the charge was heated to about 800 °C to drive

off any moisture that may have remained in the system. Then, gettered

argon was introduced to the system to a pressure of five hundred torr.

Finally, the charge was melted and held molten about twenty minutes

with the power reduced gradually until solid began to form at which time

the specimen was held at temperature about ten minutes. Power was then

turned off entirely and the specimen allowed to furnace cool.

Metal 1 ographic examination indicated that there was present less

than one percent secondary phase present in the prepared compound. The

primary phase was confirmed by x-ray examination to be Cr^C^.

- 29-
1. REFERENCES

C. S. Tedmon, Jr., D. A. Vermilyea, and J. H. Rosolowski, J. Electrochem. Soc., 118, 192 (1971).

2. R. L. Cowan, II and C. S. Tedmon, Jr., "Intergranular Corrosion of

Iron-Nickel-Chromium Alloys", General Electric Report No. 72CRD298,

Technical Information Series, October 1972.

3. A. E. Pickett, Stress Corrosion Tests of Reactor Structural Materials

in Simulated BWB Environments, General Electric Report NEDE-13179, April

1971.

4. J. Kruger and J. R. Ambrose, Technical Summary Report Number 4,

NBSIR 73-244, "The Role of Passive Film Growth Kinetics and Properties

in Stress Corrosion and Crevice Corrosion Susceptibility", 1973.

5. J. R. Ambrose and J. Kruger, Corrosion, 28, 30 (1972).

- 30 -
PART IV

[To be published in book entitled "Electrochemical Techniques for Corrosion",

NACE, Houston (1977).]

NEW APPROACHES TO THE STUDY OF LOCALIZED CORROSION

Jerome Kruger

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

A number of approaches have been developed over the years to study

localized corrosion. This paper will single out some of the newer approaches that have attempted to focus on a number of relevant issues

that must be always borne in mind if one wishes to obtain a valid understanding of local corrosion processes., This paper will not review all

the relevant issues, nor will it be an all-inclusive description of the

various new ways in which localized corrosion can be studied. Instead,

it will mostly concentrate on the results of some experimental approaches

that have had an influence on changing some of the ways localized

corrosion has been considered in the past. This paper is not intended

as a comprehensive look but rather a more narrow, personal, and, perhaps,

opinionated enumeration of important issues to be considered when studying localized corrosion.

RELEVANT ENVIRONMENTS

One of the major reasons why many studies of localized corrosion

have failed to reveal important insights into the processes controlling

such corrosion has been the failure to use relevant environments. This

is so because localized corrosion usually takes place within the restricted

confines of a crack, a pit, or a crevice. Because diffusion of ions in or out of the confines of this restricted space (called by B. F. Brown

an "occluded cell") is difficult, concentrations of these ions are built

up or decreased. Thus, the environment within a pit is not that of the

environment of the solution that exists outside the pit. Unfortunately,

it is the environment outside the pit in which many workers choose to

carry out experiments. For example, a very important reaction occurring

in an "occluded cell" is the hydrolysis reaction that can lower the pH

below that of the surrounding environment in the following way:

M

+X

+ X H

2 0 - M ( OH

)

x + X H+

. [1]

The relevant environment to be used in a study or to be considered in

the interpretation of an experiment would be the one that takes the pH

of the environment produced by reaction [1] into consideration.

- 31 -
While it has been recognized for a number of years that the en- vironment trapped in the interior of restricted regions on a metal

(cracks, pits, and crevices) has a different composition than the bulk

environment surrounding the metal, it has been only in recent years that

such knowledge was considered important, and serious attempts made to

measure the composition of the occluded environment in a crack. (An

important exception is the use of ferric chloride solutions of low pH to

carry out screening tests of the resistance of alloys to pitting.) Brown

and co-workers (1-3), using the insights provided by the large body of

work of Pourbaix (see the review in (4)) that systematized the role of

pH and potential in all corrosion processes, set about to measure these

parameters in cracks. Their original approach was to immobilize the

corrodent within stress corrosion cracks by freezing it in liquid

nitrogen, breaking open the crack mechanically, and then analyzing the

thawed out solution removed from the crack. A later development used

microelectrodes to determine the pH and potential in the crack's environment. Other developments in this fast moving field have been the use of a capillary to sample the crack solution (5) and the combination of a capillary technique with thin layer chromatography (6) to examine the

solution in crevices.

The value of these crack pit or crevice chemistry techniques

becomes evident when Fig. 1 is inspected. It shows for a large number

of steels of different compositions in environments of either high or

low pH that the potential and pH inside the crack falls in a region of

the pH potential diagram (Pourbaix diagram) where one expects both

corrosion and, because the points on Fig. 1 are below the line marked H,

the presence of hydrogen. Thus, any failure prevention measures using

bulk composition or potentials could be based on non-real istic conditions

and lead to an improper choice. A good discussion of the electrochemical

ramifications of knowing the proper environment in a crack is given by

Pourbaix (4).

RELEVANT ALLOY STRUCTURE AND COMPOSITION

In many cases of localized corrosion, differential rates of attack

for the different phases present in an alloy are observed. For example,

the difference in potential between the phase making up a grain boundary

(or grain boundary region) and the interior of a grain may be the driving

force for intergranular corrosion. A similar situation can exist when

pitting or stress corrosion is observed. Thus, it becomes important,

when one wishes to disentangle the origins of localized attack, to study

in detail the process occurring on the individual phases. A fruitful

approach is to study, for example, the electrochemistry of alloys that

are identical or similar to the relevant phases that make up the different

regions of a real alloy whose localized corrosion one seeks to understand.

Figs. 2 and 3 from some work by Ugiansky (7) show the kind of

results that can be obtained with such an approach. The anodic polarization curves carried out in a 1.0 M NaCl , pH 7, 0

2 free solution on

- 32 -
alloys simulating three important phases or regions found on Al-Zn-Mg

alloys (Fig. 2) show only minor differences between the electrochemical

behavior of the precipitate phase in the grain boundary, the precipitate

free zone (PFZ) which occurs near the grain boundary in the real alloy,

and the grain interior matrix. When 1.0 M NaNOo is substituted for 1.0

M NaCl (Fig. 3), however, the three alloys simulating the different

regions in a complete alloy exhibit markedly different electrochemical

behavior; the precipitate phase that exists in the grain boundary ex- hibits current densities (corrosion rates) that are over three orders of

magnitude higher that the other two phases examined. Thus, the approach

of studying relevant alloys (in this case simulating the different

phases or regions of an Al-Zn-Mg alloy) has yielded valuable insights

into the origins of intergranular attack of such alloys.

Not only are different phases or compositions important, but also

structural variations, such as different crystal lographic orientation of

the grains of a metal or the presence of variations, for example,

dislocations, within the same orientation, can be a factor underlying

localized attack. For example, studies by Kruger (8) on single crystal

iron surfaces showed that pitting was more predominant on surfaces whose

orientations were near the {110} orientation than on those of other

orientations (Fig. 4). This example illustrates the kind of information

one can gain by working with surfaces of known structures. By knowing

the orientation of the surface one is examining, one can determine which

are the relevant structures that affect the localized corrosion processes

whose origins are being unraveled.

RELEVANT CRITICAL POTENTIALS

There are vast numbers of experiments and discussions (9-13) in the

literature on the factors or underlying causes of the existence of a critical potential for pitting and for stress corrosion. This is so because the concept of a critical potential, E c

, below which localized

attack does not occur has immediate practical consequences. Obviously,

if one knows E

c for a given alloy and environment one then has a great

incentive to control the potential of that system so that it never goes

above E c and thereby, presumably, never suffers localized attack. One

does, however, have to know the relevant E

c

. It is because. the relevant

E c is not always known because it is measured incorrectly that the

concept of a critical potential is held in disrepute by some workers

(9). Even if one can measure a relevant E c , it is by no means established that all pitting or stress corrosion will be stopped by keeping a system below E c

. Nevertheless, one can never establish the practical

validity of the concept until it is assumed that the correct, that is,

relevant value of E c is used.

How is the relevant E c determined? The answer to this question

depends on what process or phenomenon determines the value of E c

.

- 33 -
Unfortunately, this has not been established although a number of

suggestions have been advanced (14). The different explanations of the

origin of E c can be listed as follows:

a) E c determined by potential of zero charge - Almost all mechanisms of localized corrosion 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 some mechanisms for pitting require that damaging anions adsorb and thereby

displace the ions forming the passive film, such an explanation of E would be reasonable, and the effect of such variables as the nature

c of the metal, concentration of anions, and temperature all affect Ec and pzc (see 15,16). Because it is difficult to measure the pzc for

solid electrodes (15) especially when adsorption may coincide with the

initiation of breakdown processes, there have been no good experimental

verifications that E c is related to the pzc.

b) E c is determined by film pressure to fracture film - Sato (17)

has identified E c with the critical potential above which the film

pressure exceeds the critical compressive stress of film mechanical

breakdown. The film pressure depends upon both electrostriction and

interfacial effects.

c) E c is determined by the potential of formation of an unprotective

f i 1m - One possible factor determining E c which may not apply to all cases may be that it is the potential above which an unprotective film

may form or deposit. Ambrose and Kruger (18) suggest that this may be

the case for iron.

d) E c is determined by repassivation kinetics - This concept of E c is based on kinetic considerations. Some studies (19-22) have suggested

that a competition between repassivation and breakdown of films is

important in determining initiation of localized attack. Videm (19) has

proposed that the value of E c for A1 is determined by the relative rates

of these 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 can affect repassivation kinetics

(23,24) and because, a number of workers (20-22,25,26) suggest, initiation of localized attack depends on a dynamic balance between breakdown and repair, the determination of E c by repassivation kinetics is a possibility to be considered. (More will be mentioned about repassivation

in the section on Relevant Basic Processes.)

Thus far, it is explanation d) that has led to new approaches for

determining E c

. These show some promise at yielding more relevant

values. Notable among these approaches is that of Pessel and Liu (13).

Their technique involves scratching the alloy surface whose E c is to be

- 34 -
determined at various potentials set by v a potentiostat in the environment under examination. When the potential is below E c

, the current

observed upon scratching rapidly returns to the original values observed

before scratching. When the potential is above E c

, no or very little

repassivation occurs and the current either remains near the value found

after scratching or it increases. Fig. 5 shows the kind of current

transients observed above and below E c

. A major advantage of this

approach is that it eliminates the induction time for localized corrosion

to be initiated. The usual methods for determining E

c do not have this

virtue. With the common methods one may either scan potentiodynamical ly

a potential region or potentiostatically sample a series of potentials

in succession staying at each potential for varying amounts of time.

When a value for the potential is reached at which the current measured

starts to increase, this value is identified as E c

. However, it is

known that there exists an induction time for localized corrosion

attack. Therefore, unless one stays at all potentials for a time

greater than the induction time, one can never be assured that a lower

potential would not have been E_. Since the induction time can be very

long depending on the history of the sample studied (see Ambrose and

Kruger (l8))> it is no wonder that the E

c determined by the usual

methods has been considered by some to be of questionable value (9).

Because Pessel and Liu's technique eliminates the problem of time of

initiation, it appears that the E c determined by its use is a more

relevant value.

Likewise, the potential identified by Pourbaix (11) as the protection

potential, Eg* may also be more relevant. This potential, which is more

active than E

c

, is the potential below which localized corrosion once started can be stopped. Pessel and Liu's work indicates that when E c is determined by their technique, a way that eliminates the time for

initiation of attack to occur, E

c = Ep.

RELEVANT BASIC PROCESSES

In the preceding section a technique for determining E c was considered

to give relevant values because it was based on the supposition that

repassivation was a process strongly affecting localized attack. A

number of workers (19-21) have suggested that the rates of two major

basic processes, breakdown of passivity and repassivation, are of great

importance in determining whether localized attack will occur. By

focusing on these two basic processes, breakdown and repairs, some valuable approaches have been developed for studying localized cor- rosion.

Actually a more relevant environment would be that environment one would expect to develop in a pit, crack, or crevice when the surface is

exposed to the ambient environment.

- 35 -
One technique that has been suggested as actually looking at

breakdown-repair events is that concerned with examining the small

fluctuations of potential or current occurring on a corroding surface,

"noise". Iverson (27) showed that the frequency and intensity of such

noise could be markedly diminished by the addition of an inhibitor to a corroding system (Fig. 6). Videm (19) and Okamoto et al . (28,29) have

suggested that the noise represents breakdown and repassivation events

and that the dynamic competition between these two basic processes

determines whether localized attack such as pitting will be initiated.

The function of the chloride ion in pitting is to tilt the scale towards

breakdown and corrosion by slowing down or stopping effective repassivation (19,21).

If repassivation kinetics is, therefore, an important basic process,

attention must be focussed on it and ways developed to separate its

measurement out from the breakdown process that always occurs along with

it during localized corrosion. Scully (30,31) was one of the earliest

workers to emphasize the importance of repassivation kinetics, and it

has become an important element, especially in the film rupture metaldissolution mechanism of stress corrosion (SCC), because combined with

the concept of slip step dissolution, it explains why restricted lateral

dissolution occurs - the slower the repassivation, the more of the metal

slip step that has emerged and ruptured the film can be dissolved before

new film forms. If too little dissolves, no SCC occurs; if too much

dissolves, the dissolved areas are not restricted enough, and the crack

is too blunt to propagate. Only when the environment and alloy provide

a system where the repassivation kinetics produce a highly restricted

region for dissolution is SCC found. In addition to Scully, Staehle

(32) has given a detailed discussion of the significance of repassivation

kinetics. \

These considerations have resulted in a number of techniques aimed

at determining the repassivation kinetics (33-38). These techniques

produce a bare surface either by straining an alloy wire or by scratching or abrading a surface. In most of these techniques, the specimens

are either strained or scratched while maintaining their potentials at a fixed value and measuring the current transient that takes place during

the dissolution-repassivation event that occurs subsequent to exposure

of bare metal. The rate of current decay is related to the rate of

repassivation. A new and useful technique for measuring susceptibility

to SCC by looking at repassivation kinetics is that developed by Galvele

et al. (39). This technique determines anodic and cathodic polarization

curves for both an unstrained and strained electrode. The ratio between

the corrosion current for the unstrained electrode to that of the

strained can then be used to assess susceptibility.

In such mainly electrochemical techniques, one cannot separate out

the current involved in dissolution from that involved in passive film

formation. Using a technique developed by Ambrose and Kruger (38),

however, it is possible to measure separately the film regrowth transient

- 36 -
ij by optical means (el 1 ipsometry ) and to compare it to the overall dissolution-

: repassivation current. With this technique, called tribo-ell ipsometry,

j one removes the protective film on a metal in an environment by means of

! a polishing wheel and then records both the film growth kinetics and the

j current as shown in Fig. 7. In Fig. 7 we can see that the rate of film

! growth on a titanium alloy is considerably slower in the susceptible

| chloride environment than it is in the nonsusceptible nitrate solution.

; With tri bo-el 1 ipsometry one can determine the ratio of the total current

to that involved in repassivation only. This ratio R

p * when the exposed

surface area is known is given by the expression

Q

d V

‘ 1 + q;

[ 2 ]

where = charge involved in dissolution and Q r

= charge involved in

repassivation.

Rp* is a measure of the effectiveness of the repassivation process.

Work on the cracking of low carbon steel (40) has shown that SCC does

not occur when the effectiveness of repassivation is either too high or

too low. Thus, for low carbon steel, at the time required to form a complete film, R

p * is 2.8 in 1 N NaN02 (25 °C) where it is nonsusceptible;

R

p * is 26 in 4 N NaN03 (90 °C) where it does undergo SCC, and R

p * is

> 75 in 1 N NaCl (25 °C). where it undergoes widespread pitting attack

but not SCC. Thus, a delicate balance between repassivation and dissolution must be maintained to lead to SCC. The determination of this

balance is dependent on the environment, the alloy, and finally the

metallurgy and mechanics of the system that control the rate at which

bare metal sites (generally slip steps) are produced as well as their

density and size.

It is on the rate at which bare metal sites are produced that

another new approach to localized corrosion (primarily SCC) is based.

This approach uses the constant strain rate technique which gives some information on the beneficial breakdown processes occurring in SCC.

This technique, which produces experimental data as depicted in Fig. 8, can be applied both to mechanistic studies as well as used for routine

testing. The mechanistic importance of the technique is that it ties

the metallurgically controlled rate of bare metal exposure, resulting

from film rupture by slip step emergence, to the chemical process of

repassivation of the exposed metal. This can be seen by examining the

significance of the three regions shown in Fig. 8. For Region A where

strain rates are too low to lead to SCC, the rate of repassivation is

high enough to prevent any significant and damaging interaction between

the environment and the bare metal exposed during straining. At the

opposite extreme, Region C, the strain rate is so rapid that the interaction of the exposed metal with the environment is of no consequence

because the rate that the metal fails by ductile fracture exceeds the

rate by which the environment can affect fracture via any of the possible SCC mechanisms (dissolution, hydrogen embrittlement, or adsorption

of damaging species). It is only in the intermediate range of strain

rates, Region B, where the rate of production of bare metal sites has a

- 37 -
value sufficiently high so that the rate of repassivation does not

prevent environmental access. However, the strain rate is sufficiently

low so that environmental interaction and not ductile pulling apart

produces fracture. Under these conditions SCC takes place. If SCC is

not possible for the system studied, the dip shown in Region B will not occur.

Does this test, whose development and increasing use was stimulated

for mechanistic reasons and was mainly guided by the implications of

film rupture and repassivation, provide a valid routine material -environmental

evaluation test? Parkins (41 ) has pointed out that there are certain

features of the constant strain rate technique that make it a useful

routine laboratory test. These features are as follows:

(1) It is a relatively severe test so that it will promote laboratory SCC failures where other tests on smooth specimens will not

unless inordinately long testing times are used;

(2) It always produces fracture by SCC or some other mechanism and

therefore is a positive test;

(3) The time of testing is relatively short.

These three somewhat interrelated attributes, severity, positiveness,

and rapidity, are especially valuable in comparative testing where one wishes to intercompare materials and environments.

The promise that the constant strain rate approach offers for

laboratory testing is somewhat tempered by the lack of good ways to

quantify the results of the tests. For example, how does one relate the

strain rate in Region B where SCC susceptibility is the greatest to

crack growth rate? Moreover, it is not known what relation, if any,

this strain rate bears to the strain rate that exists at the tip of a growing crack. Parkins (41) has made some valuable qualitative connections

between this value and the threshold stress for SCC or the threshold

value of the stress intensity factor (an important parameter in the

fracture mechanics approach to SCC testing (42)) but the quantitative

relationships require new theoretical and experimental advances.

Despite these problems, the constant-strain rate tests have provided

evidence of SCC in systems where other tests have been difficult or not

possible to apply (43-46).

RELEVANT FILMS

The protective films that form on alloy surfaces in a given environment are of crucial importance in controlling localized attack because

their ability to withstand the breakdown processes determines whether

such attack will occur. In attempting to determine how these films

affect localized attack one must be sure that certain relevant con- siderations with respect to the film's ability to affect localized

- 38 -
corrosion are borne in mind. Three such important, but not all-inclusive,

considerations will be suggested here, 1) pretreatment, 2) structure,

and 3) origin.

The first consideration, pretreatment, is concerned with the history

of the film prior to its exposure to an environment that causes localized

attack - generally one containing aggressive anions. It is because this

consideration is ignored that potentiodynamic scans lead to erroneous

values of E c , as discussed in an earlier section. This is so because

pretreatment can markedly affect the time required to initiate breakdown. For example, Ambrose and Kruger (18) 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 that the defects in the film, through which

penetration must proceed, were removed by the annealing. Similar

results for Ta were found by Vermilyea (47). Fig. 9 shows, by using the

technique of ell ipsometric spectroscopy (48),- that the annealed films on iron exhibit entirely different optical changes when exposed to Cl"

ions. It is thus obvious that the relevant film to be considered in a study of localized corrosion is the one whose pretreatment history is

known.

Pretreatment by annealing does not always increase the time for

localized corrosion initiation to occur. Unlike the situation for iron,

as just described, Okamoto (22) found that annealing the passive film on stainless steel shortened the induction time for the initiation of

pitting. The reason for this entirely opposite behavior for stainless

steel as compared to iron lies in the second consideration, film structure.

The passive film on stainless steel is noncrystalline or "glassy" (22,49)

and such a glass-like structure provides few diffusion paths for penetration by aggressive anions. When annealed, however, the noncrystalline structure of the passive film on stainless steel can crystallize

and the new grain boundaries formed offer greater opportunities for

diffusion. It can thus be seen that one has to consider the structure

of the relevant passive film in studying localized corrosion.

The influence of alloy composition on structure and composition of

the passive films is intimately tied in with the third consideration,

the origin of the film; or more precisely, does the relevant film being

considered come from growth on the alloy or from precipitation or formation

from the environment? To illustrate the point that is being made,

consider the problem that although Mo in a stainless steel is known to

be beneficial in preventing or retarding pitting or crevice corrosion,

recent work (50,51) has shown that it is not present in the passive

films formed in such alloys. Its beneficial effect can only come from

its effect on propagation rather than on initiation where the initial

passive film properties are important. Streicher (52) has shown this to

be so and recent repassivation kinetic studies by Kodama and Ambrose

(53) found that Mo may be involved in inhibiting pit propagation by

precipitating a protective layer of ferrous molybdate from solution, the

- 39 -
molybdate being formed when some of the alloy goes into solution during

pit initiation. Thus, the relevant protective film affecting localized

corrosion is not the original film formed by passivation of the alloy

surface but a film whose origin is a precipitation reaction occurring in

solution between ions that were formed during the initiation stage.

CONCLUDING REMARK

I A fruitful study of localized corrosion requires the same weeding

out exercise required in watching some modern or avant garde plays.

Unless one separates out the relevant characters (relevant environment,

alloys, potentials, films) from those that do not matter and follows the

relevant plot lines (relevant basic processes), one will not get much

out of the play, nor will one get much understanding of the localized

corrosion process being studied.

ACKNOWLEDGMENT

I am grateful to the Office of Naval Research which supported this

work under contract NAONR 18-69 NRO 36-082.

- 40-
REFERENCES

1. B. F. Brown, C. T. Fuji i , and E. P. Dahlberg, J. Electrochem. Soc., 1 1

6

218 (1969).

2. G. Sandoz, C. T. Fujii, and B. F. Brown, Corrosion Sci., H) 839 (1970).

3. J. A. Smith, M. H. Peterson, and B. F. Brown, Corrosion, 26, 539 (1970).

4. M. Pourbaix, Theory of Stress Corrosion Cracking of Alloys, J. C. Scully,

Editor, NATO, Brussels, p. 17 (1971).

5. H. Leidheiser, Jr., and R. Kissinger, Corrosion, 28, 218 (1972).

6. F. D. Bogar and C. T. Fujii, Naval Research Laboratory Report 7690,

March, 1974.

7. G. M. Ugiansky, Dissertation, Ohio State University, 1976, to be published.

8. J. Kruger, J. Electrochem. Soc., 1 06 , 736 (1959).

9. M. J. Pryor, Localized Corrosion, Staehle, Brown, Kruger, and Agrawal

,

Editors, NACE, Houston, p. 2 (1974).

10. B. E. Wilde, Localized Corrosion, Staehle, Brown, Kruger, and Agrawal,

Editors, NACE, Houston, p. 342 (1974).

11. M. Pourbaix, Localized Corrosion, Staehle, Brown, Kruger, and Agrawal,

Editors, National Association of Corrosion Engineers, p. 12 (1974).

12. S. Smialowska and M. Czachor, Localized Corrosion, Staehle, Brown, Kruger,

and Agrawal, Editors, NACE, Houston, p. 353 (1974).

13. N. Pessall and C. Liu, Electrochemica Acta, lj>, 1 987 (1 971 ).

14. J. Kruger, Proc. Passivity and its Breakdown on Iron and Iron Base Alloys,

R. W. Staehle and H. Okada, Editors, NACE, Houston, p. 91 (1976).

15. J. 0'M. Bockris and A. K. N. Reddy, Modern El ectrochemi stry , Vol . 2,

Plenum Press, New York, p. 706 (1970).

16. S. D. Argade and E. Gileadi, Electrodeposition, E. Gileadi, Editor,

Plenum Press, New York, p. 87 (1967).

- 41 -
17. N. Sato, Electrochemica Acta, 19_, 1683 (1971 ).

18. J. R. Ambrose and J. Kruger, Proc. 4th International Congress on

Metallic Corrosion, NACE, Houston, p. 698 (1972).

19. K. Videm, Kjeller Report KR-149, Institutt for Atomenergi , Kjeller,

Norway (1974).

20. J. Kruger and J. R. Ambrose, NBS Report NBSIR 74-583, September (1974).

21. J. Zahavi and M. Metzger, Localized Corrosion, Staehle, Brown, Kruger,

and Agrawal , Editors, NACE, Houston, p. 547 (1974).

22. G. Okamoto, Corrosion Science, T3, 471 (1 973).

23. J. R. Ambrose and J. Kruger, Proc. 5th International Congress on Metallic

Corrosion, NACE, Houston, p. 406.

24. J. R. Ambrose and J. Kruger, J. Electrochem. Soc. , 1 21 , 599 (1974).

25. J. R. Galvele, S. M. deMicheli, I. L. Muller, S. B. deWexier, and I. L.

Alanis, Localized Corrosion, Staehle, Brown, Kruger, and Agrawal, Editors,

NACE, Houston, p. 580 (1974).

26. K. J. Vetter and H. H. Strehblow, Localized Corrosion, Staehle, Brown,

Kruger, and Agrawal, Editors, NACE, Houston, p. 240 (1974).

27. W. P. Iverson, J. Electrochem. Soc., 115 61 7 (1968).

28. G. Okamoto, T. Sugita, S. Nishiyama, and K. Tachibana, Boshoku Gijustou,

23, 439 (1974).

29. G. Okamoto, K. Tachibana, S. Nishiyama, and T. Sugita, Passivity and

its Breakdown on Iron and Iron Base Alloys, R. W. Staehle and H. Okada, Editors,

NACE, Houston, p. 106 (1976).

30. J. C. Scully, Corrosion Sci . , 7_, 197 (1 967).

31. J. C. Scully, Corrosion Sci., 8, 51 3 (1968).

32. R. W. Staehle, Theory of Stress Corrosion Cracking in Alloys, J. C. Scully,

Editor, NATO, Brussels, p. 223 (1971).

- 42 -
33. T. P. Hoar and J. M. West, Proc. Roy. Soc., 268A , 304 (1962).

34. H. Leidheiser, Jr., and E. Kellerman, Corrosion, 26, 99 (1970).

35. T. Murata and R. W. Staehle, Proc. 5th Inti. Cong, on Metallic Corrosion,

NACE, Houston, (1974).

36. N. 13. Tomashov and L. P. Vershinina, Electrochimica Acta, 1_5, 501 (1970).

37. J. R. Ambrose and J. Kruger, Corrosion, 28, 30 (1972).

38. R. B. Diegle and D. A. Vermilyea, J. Electrochem. Soc., 1 22 , 180 (1975).

39. J. R. Galvele, S. B. deWexler, and I. Gardiazabal, Corrosion, 31_, 352

(1975).

40. J. R. Ambrose and J. Kruger, Proc. 5th Inti. Cong, on Metallic Corrosion, 1 l Tokyo, NACE, Houston, p. 406 (1974).'

41. R. N. Parkins, Theory of Stress Corrosion Cracking in Alloys, J. C. Scully,

Editor, NATO, Brussels, p. 449 (1971).

42. B. F. Brown, Ed., Stress-Corrosion Cracking in High Strength Steels and

in Titanium and Aluminum Alloys, Naval Research Laboratory, U.S. Government

Printing Office, Washington, 1972.

43. W. R. Wearmouth, G. P. Dean, and R. N. Parkins, Corrosion, _29, 251 (1 973).

44. R. N. Parkins, F. Mazza, J. J. Royuela, and J. C. Scully, Brit. Corr.

J., 7, 154 (1972).

45. J. C. Scully and D. T. Powell, Corrosion Sci., ]_0, 371 (1970).

46. D. C. Deegan and B. E. Wilde, Corrosion, 2J3, 310 (1 973).

47. D. A. Vermilyea, J. Electrochem. Soc., 104 , 485 (1957).

48. C. L. McBee and J. Kruger, Localized Corrosion, Staehle, Brown, Kruger,

and Agrawal , Editors, NACE, Houston, p. 252 (1974).

49. C. L. McBee and J. Kruger, Electrochemica Acta, T_7 , 1 337 (1972).

50. J. B. Lumsden and R. W. Staehle, Scripta Metallo, 6, 1205 (1972).

51. C. L. McBee and J. Kruger, Passivity and its Breakdown on Iron and

Iron Base Alloys, R. W. Staehle and H. Okada, Editors, NACE, Houston, p. 131

(1976).

- 43 -
52. M. A. Streicher, J. Electrochem. Soc., 103, 375 (1956).

53. T. Kodama and J. R. Ambrose, NBS Report NBSIR 75-916, October (1975).

Submitted to Corrosion.

54. B. F. Brown, Theory of Stress Corrosion Cracking in Alloys, J. C. Scully,

Editor, NATO, Brussels, p. 197 (1971).

- 44-
Fig. 1. ential-pH diagram showing the potential I

H 'val u<;S

ck tip environments for a number of high strength

im Brown (54).

in the

steels.

-45-
1.0M NaCI, pH 7, 02 FREE

Fig. 2. Polarization curves obtained in a 1.0 M NaCI solution, pH 7

O2 free for different alloys simulating the various regions

or phases existing in an Al-Zn-Mg alloy. From Ugiansky (7)

-46-
1.0M NaN03

, pH 7. 02 FREE

Fig. 3. Polarization curves obtained in a 1.0 M NaN03 solution, pH 7,

O2 free for different alloys simulating the various regions

or phases existing in an Al-Zn-Mg alloy. From Ugiansky (7).

-47 -
Fig. 4. Stereographic triangle showing pitting on grains of iron

having different crystallographic orientations. From

Kruger (8).

-48-
KT4 -

Above E

10-5

Scratch

Below E c

Repossivotion

Breakdown

K>'6| I I I I I I I I J «J 0 1 2 349678910

Time , minutes

Fig. 5. Scratch-test current-time curves for specimen held

potentiostatically above Ec (breakdown) and below E c (repassivation).

-49-
Fig. 6. Periodic potential fluctuations on A--a mild steel surface in

0.1% NaCl solution and B--after 1% NaN02 has been added to the

NaCl solution. Note how the inhibitor NaN02 reduces the pulse

intensity. From Iverson (27).

- 50-
Fig. 7. A comparison of oscilloscope repassivation transients for Ti-8

Al-1 Mo-1 V alloy in 1 N MaCl and NaNC>3 solutions at +706 mV (SHE)

(el 1 ipsometric film growth-jagged trace; current-solid trace).

From Ambrose and Kruger (23).

- 51 -
i

% Reduction in Cross Sectional

Area

Fig. 8. A schematic representation of the results obtained in constant

strain-rate tests. Region A—repassivation is rapid enough

and strain rate slow enough to prevent damaging interaction

with the environment. Region B—strain rate high enough so that repassivation does not prevent environmental access. Region C—strain rate is so rapid that ductible pulling apart

takes place.

- 52-
I

WAVELENGTH, XilO*

Fig. 9. A comparison of the effect of the addition of chloride ions on the el 1 ipsometric spectrum (relative phase retardation,

A, vs. wavelength) obtained for a passive film formed on iron

in sodium tetraborate-boric acid solution by polarizing at

1000 mV (SHE) to the spectrum of a similar passive film

annealed for 1 hr. at 75°C. a) unannealed film; b) annealed

film. From McBee and Kruger (48).

-53-
NBS-1 14A (REV. 7-73)

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBSIR 76-1170

2. Gov’t Accession

No. I 3. Recipient’s Accession No.

4. TITLE AND SUBTITLE

The Role of Passive Film Growth Kinetics and Properties in

Stress Corrosion and Crevice Corrosion Susceptibility - Technical Report No. 7

5. Publication Date November 1976

6 . Performing Organization Code

7. AUTHOR(S)

J. Kruger and J.R. Ambrose

8 . Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 3120448

11. Contract/Grant No.

'I AONR 18-69, NR036-082

12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

Office of Naval Research , Department of the Navy

Arlington, Virginia 22217

13. Type of Report & Period

Covered

interim 10/75-11/76

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200- word or less bibliography or literature survey, This report consists of

Fe-Mo alloys on crevice

repassivation in a crevi

repassivation rates in a the composition of metal

steel. These alloys are repassivation kinetics,

rosion.

factual summary of most significant information. If document includes a significant

mention it here.)

four parts as follows: I-A study of the influence of Mo in

corrosion. It was found that >5% Mo is needed to affect

ce. II-The description of a new technique for measuring

crevice. I I I-A description of alloys prepared to simulate

near grain boundaries of sensitized austenitic stainless

used for studies of the effect of sensitization on

IV-A review of new approaches in the study of localized cor17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons)

Crevice corrosion; Iron-molybdenum alloys; Localized corrosion; Repassivation; Sensitization; Stainless steel

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