NBSIR 74-583

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

September 1974

Technical Summary Report Number 5

Prepared for Office of Naval Research

Department of the Navy

Arlirjgton. Va. 22217

NBSIR-583

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

September 1974

Technical Summary Report Number 5

DISTRIBUTION OF THIS DOCUMENT IS UNLIMITED

Prepared for Off ice of Naval Research

Department of the Navy

Arlington, Va. 22217

U. S. DEPARTMENT OF COMMERCE, Frederick B. Dent. Secretary

NATIONAL BUREAU OF STANDARDS, Richard W. Roberts. Director
ii! •..1
PART I

To be submitted for publication in Corrosion

Effect of Molybdenum Alloying

on the Repassivation and Depassivation

of Ferritic Stainless Steels

by

J. R. Ambrose and J. Kruger

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

ABSTRACT

A study of the effect of alloying chromium and molybdenum in

ferritic stainless steels has been made using techniques to study

repassivation kinetics (tribo-el 1 ipsometry) and depassivation

kinetics (a recently developed ell ipsometry technique for studying

optical changes occurring within a crevice). Results indicate

that chromium additions affect both reoassi vation and depassivation

processes v/hile molybdenum appears to affect only the repassivation

processes. A discussion is given of a generalized mechanism for

localized corrosion which postulates that susceotibil i ty to attack

is determined by the competititon between depassivation and repassivation processes.

-1-
Effect of Molybdenum Alloying

on the Repassi vation and Depassivation

of Fern" tic Stainless Steels

by

J. R. Ambrose and .J. Kruger

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

Although additions of molybdenum to ferritic stainless steels

have been known for some time to enhance resistance to pitting

and crevice corrosion (1,2), attempts to determine the origins of

this increased resistance have, for the most part, been limited

to electrochemical measurements of corrosion potentials and

anodic currents under applied potential on various alloys.

Molybdenum concentration in the alloy was found to be directly

proportional to the corrosion and the critical pitting potential,

and inversely proportional to the peak current density required

for passivation (3). What these results do not show is how these

variations in electrochemical behavior are affected by the

presence of molybdenum in the metal. It would seem that if corrosion resistance were attributable to the presence of a protective

film, then perhaps incorporation of molybdenum in the film is a factor, but Auger spectroscopy measurements actually reveal a surface depletion of the element (4). If molybdenum is indeed not

in the passive film, its ability to impart to an alloy a greater

resistance to localized corrosion must tie in with a mechanism

more subtle in origin than the production of more protective film.

This paper explores the role that molybdenum plays in concert

with the chromium in ferritic stainless steels in the dynamic interplay of two processes proposed to control localized corrosion:

depassivation and repassi vation . The broader term "depassivation"

is used rather than "breakdown" because it includes any process

that removes a protective film from a surface. Thus, film rupture

under the influence of external stress, which leads in some cases

to stress corrosion, would be covered by the more inclusive term.

EXPERIMENTAL

The materials used in this study were a series of Fe-Cr-Mo

alloys with approximately 14, 17, and 25 percent chromium and 0, 2,

and 5 percent molybdenum by weight. Compositions of all alloys are given in Table I. One cm square specimens were cut from the received

material and cemented to aluminum plates with silver epoxy. The

-2-
assembly was mounted in clear lucite, then drilled and tapped through

the aluminum for electrode assembly. Specimens were polished on silicon carbide papers and rotary polished with 1 ym diamond paste.

Specimens were degreased with spectro grade methanol and dried in a jet of compressed air. All experiments were performed in 1.0 N NaCl

solution prepared from reagent grade chemical and distilled water

(7 X 10-7 ohm-l cm-l)

.

The two experimental techniques employed in this study were

selected in order to measure the rates of the two competing processes

affecting corrosion resistance.

Repassi vation Kinetics

Repassivation kinetics were studied using tri bo-el lipsometry,

the apparatus and techniques of which have been described elsewhere

(5,6). Here depassi vation is caused by in situ mechanical abrasion,

and film growth rates are measured with an ellipsometer while

potentiostatic current transients are simultaneously being recorded.

Although the limiting film thickness on this series of ferritic

stainless steels can, depending upon applied potential, exceed 1 nm,

the film that forms during the repassivation transient and which is

responsible for reducing the anodic current density following film

removal to less than 50% of its maximum value, is on the order of

0.1 nm and forms quite rapidly, in less than 15 ms. Since no ellipsometer intensity measurements were obtainable in the early

stages of film growth, the oscilloscope scale was expanded to include

current and ellipsometer parameter transients up to 100 ms. However,

within this range, changes in film thickness were small, barely

detectable above background photomultiplier noise levels, and could

not be associated with molybdenum content of the alloys studied

(Fig. 1). The rate of current decay, on the other hand, was dependent upon alloy composition for all alloys studied, and it is

this result that will be presented here. The ellipsometer photo

detection system is being modified to include a thermo-cooled gallium

arsenide photomultiplier tube which is reported to reduce background

noise by some three orders of magnitude. This system, when combined

with a high voltage tungsten filament light source, might offer

sufficient sensitivity to measure film thickness variations during

this time interval.

Depassi vation Studies

Depassi vation presumed to occur during crevice corrosion was studied using a new technique also utilizing the ellipsometer (7).

Using a glass plate to create a crevice, measurements of optical

changes occurring during breakdown could be made, and these results

correlated with electrochemical measurements.

-3-
Table I.

Alloy Composition

Cr Mo C N Ni Si P S

13.68 0.08 0.003 0.004 NA NA NA NA

13.60 1.88 0.002 0.004 NA NA . NA NA

17.35 <0.01 0.002 0.003 NA NA NA NA

17.61 2.02 0.004 0.004 NA NA NA NA

17.38 4.75 0.004 0.004 NA NA NA NA

25.66 0.002 0.002 B 0.13 B B B 25.26 2.05 0.002 0.017 0.13 B B B 25.25 5.24 0.002 0.013 0.13 0.19 0.12 0.002
time

^' ^' P???ncnIf/^^'^^M^"^^-'°" ^'"^"^ measurement using tri e ipsometry. Majority of film growth complete before ellipsometer measurement made (10-15 ms).
Crevice Electrode Schematic

Specimen Electrode Assembly

Fig. 2. Schematic of crevice assembly used in depassivation kinetics

studies.

-6-
The crevice assembly is the same as described earlier (7) with a quartz glass plate substituted for the pyrex (Fig. 2). All crevices

were 10"^- cm wide. The same alloys and solutions were used in this

phase as in the repassi vation kinetics studies. Each specimen was potentiostatted at its previously determined open circuit corrosion

potential immediately after creation of the crevice (7). El 1 ipsometric

and current measurements were then made over about a 1000 minute time

span.

RESULTS

Repassi vation Kinetics

Log log plots of anodic current versus time following cessation

of abrasion for the Fe-14 Cr with 0 and 2 percent molybdenum are shown in Fig. 3. The linear nature of the plot suggests the

relationship:

i = At""" [1]

where k is the slope of the decay plot and a function of, among

other things, the molybdenum concentration in the alloy. The quantity

A in equation [1] is related to the maximum current density for

exposed bare metal and is more or less dependent upon the abrasion

efficiency (amount of surface area exposed by the carbide paper) of

the tri bo-el lipsometer system. Since the maximum current level was

not particularly reproducible, attention will not be placed on it.

The k values, however, were extremely reproducible, varying by no more than +1%. The experimental results shown here are based on the second abrasion cycle of a fresh carbide paper and A values

for different alloys can be compared on an empirical basis. From

Fig. 3 it can be seen that the addition of 2% molybdenum to the Fe-14

Cr slightly decreases the rate of current decay. The effect is much

more pronounced in the higher chromium containing alloys Fe-17 Cr

(Fig. 4) and Fe-25 Cr (Fig. 5). The values of k varied from 0.42 for

the Fe-14 Cr alloy to about 1.2 for the Fe-25 Cr-2 Mo alloy. Values

of k plotted as a function of molybdenum concentration are shown

in Fig. 6. It is interesting to note that additions of molybdenum

raise k to a limiting value quite near that found by Nagayama and

Cohen (8) for iron passivation in borate buffer solution, whose

current efficiency for film formation was essentially 100%. Secondly,

the fact that k is linearly dependent upon molybdenum concentrations

is consistent with the results of Rockel (3) who, for similar alloys

in H2SO4 solution, also found a linear dependence for critical peak

current density on percent molybdenum in the alloy. The slope of this

linear relationship is clearly a function of chromium content (Fig. 7).

When the potential of the surface was changed, the Fe-Cr alloys

containing no molybdenum showed no variation in k with an increase in

potential (Fig. 8). When molybdenum was added, however. Fig. 9 shows

that k increased with an increase in potential.

-7-
TIME , msec

3. log-ir^ i versus loq-i^ t during repassivation for Fe-14 Cr

(-045f mV SHE) and Fe-14 Cr-2 Mo (-008 mV SHE) in l.ON NaCl

solution.
10 20 30 40 50 60 70 80 90KX}

TIME , msec

Fig. 4. logiQ 1 versus loqiQ t during repassivation for Fe-17 Cr

(-049 mV SHE), Fe-U Cr-2 Mo (-062 mV SHE) and Fe-17 Cr5 Mo (+020 mV SHE) in 1 .ON NaCl solution.

-9-
<

E

UJ

q:QCDO

V

Oj^|.25Cr.2ii

I I I I 10 20 30 40 50 60 70 80 90 100

TIME, msec

Fig. 5. log-jQ i versus loa-.^ t during repassivation for Fe-25 Cr

(-044 mV SHE), Cr-2 Mo (+142 mV SHE), and Fe-25 Cr5 Mo (+042 mv SHE) in 1 .ON NaCl solution.

-10-
Fe-25Cr

Fe-l7Cr

Fig. 6. Plot of k, Aiogin t» versus weight per cent Molybdenum for

Fe-14 Cr, Fe-17 Cr, and Fe-25 Cr ferritic stainless steels,

-11-
Fig. 7. Plot of Ak/A% Mo versus chromium concentration.
Fe-l4Cr

' ' '

' 10 20 30 40 50 60 70 80 90 100

Time, msec

Fig. 8. Effect of increased applied potential on the current decay

kinetics of Fe-14 Cr. Potentials are versus SHE.

-13-
10

<

E 4

CC 3

crO

10

Fe- l4Cr-2Mo -8mV

20 30 40 50 60 70 80 90 100

TIME, msec

M^l^i-cfoi^Fp'^f? °" c^^^ent decay Kinetics of Fe-14 Cr-2 Mo. Potentials are versus SHE.

-1)4-
Thus, one can increase the rate of repassivation of an Fe-Cr

alloy either by adding more molybdenum or, for a fixed amount of

molybdenum, by raising the potential.

A relationship between repassivation kinetics and stress cor- rosion cracking susceptibility was examined by comparing a ferritic

stainless steel of composition Fe-28 Cr-4 Mo to the same alloy with

a 2% nickel addition. The addition of the nickel makes the alloy

more susceptible to stress corrosion (9). Current decay transients were measured in 1.0 N NaCl and are shown in Fig. 10. The less

susceptible material repassivates about 20% faster then the susceptible

nickel containing alloy.

Depassi vation i"" Previous studies (7) have shown that the relative phase

retardation, a, is a parameter that is sensitive to the processes

leading to film breakdown and crevice corrosion. The changes in

A with time are plotted for the^ two Fe-14 Cr alloys in Fig. 11.

The rate of decrease in A was low'ered by addition of molybdenum.

Current increased steadily du^ng the course of the experiment at

about the same rate for rboth* materials and is most likely due to

oxygen depletion within t|ie crevice. Although the three stages of

change in a (decrease, ihbrease, decrease) previously found for 304

stainless steel (7) were also observed here, it is the influence of

alloying on the first stage that is probably the more important,

since depassi vation is initiated during this period by what are generally thoughit to be increases in concentration of soluble cor- rosion products within the crevice (1). Therefore,' later stages in

the crevice, corrosion process are not included here.

The breakdown process took longer to initiate in the Fe-17 Cr

series, but the results were similar to the Fe-14 Cr alloys (Fig. 12).

A was not found to decrease to a measurable extent in the Fe-25 Cr

series of alloys.

Bates (10) recently showed that application of cathodic protection

for short intervals of time would inhibit the crevice corrosion process.

To see if the optical changes leading to breakdown were also arrested,

cathodic protection was applied and the results shown in Fig. 13.

Upon switching the applied potential some 240 mV in the cathodic

direction after the breakdown process had begun, it was possible to

arrest both a and i|; values, but the system did not return to its

original state.

DISCUSSION

Before examining the significance of the results described, it is

necessary to discuss the possible interplay of depassivation and

repassivation in localized corrosion. The concepts developed will then

be applied to an examination into the effects of adding chromium and

molybdenum to ferritic stainless steels.

-15-
TIME , msec

Fia. 10. logig i versus loaiQ t during repassivation for Fe-25 Cr-4 Mo

(-0(^3 mV SHE) and Fe-28 Cr-4 Mo-2 Ni (+094 mV SHE).

-16-
8A

2

-

Fia. 11. Chanae in relative phase retardation. A, with Ioqiq t for

Fe-14 Cr (-049 mV SHE) and Fe-14 Cr-2 Mo, (-008 mV SHE)

at within a crevice.

-17-
12. Change in relative phase retardation, ^' versus loa^Qt for

Fe-17 Cr (-049 mV SHE). Fe-17 Cr-2 Mo (-062 mV SHE), and

Fe-17 Cr-5 Mo (+020 mV SHE) in l.ON NaCl solution within

a crevice,
Potential " -18mY

Fe 17Cr

Potential to -2&8niV SHE

10 100

TIME . min

Fig. 13. Effect of application of cathodic protection on el lipsometer

parameters A and At approximately 100 minutes, after

A and \i) had begun to change due to onset of crevice attack,

potential shift to -258 mV SHE.
Depassi vation and Repassivation

If all localized corrosion is regarded as resulting from some disruption of a protective film, then those events which control access of unprotected metal to the environment will determine

susceptibility to a given mode of corrosion attack, as well as affect its rate. This generalized film disruption process will

be loosely defined as "depassi vation" . This process can result

from the application of external stress that could lead to stress

corrosion or it can result from breakdown of the film by some sort

of interaction with the environment which could lead to crevice

or pitting corrosion. Subsequent events through which the system

can be returned to its former state are classified under "repassivation".

The premise of this discussion will be that it is the balance

between depassivation and repassivation that will be responsible

for controlling the performance of a material exposed to a particular

environment.

It would be useful to know the nature of the steady state

conditions which exist at the metal solution interface. Okamoto has

suggested that a metal capable of suffering localized corrosion

experiences a dynamic process of continued breakdown and repair (17).

One possibility that is consistent with all our experimental results

is that the depassivation events in this dynamic balance occur mechanically through rupture of the surface film due to internal

straining.

Sato has recently suggested that high fields could lead to

mechanical rupture of thin films by electrostriction pressures

exceeding the compressive fracture strengths of the film (18). The

so-called film pressure this generates is given by:

P Po 8tt L where p = film pressure

p^ = atmospheric pressure

e = film dielectric constant

E = electric field

a = surface tension

L = film thickness

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

promoting corrosion attack increase film pressure by decreasing the

surface tension term in equation [2] through adsorption. Yet, this

should occur with all anions, not just a select few. Zero charge

potentials for many anions, aggressive and non-aggressive alike, are similar on mercury (23). Since zero charge potentials would also be

expected to be similar on most oxides, any increase in film electrostriction pressure due to a decrease in surface tension because of

adsorption would be a general effect and could not explain a specificity of those ions which are known to cause pitting. Increase

in film dielectric constant through a kind of anion doping would also

-20-
lead to increased film pressure and would be somewhat specific in

that ion size might be important. However, chloride ion, an aggressive species, has about the same polarizabi 1 i ty as oxide ions

and only slight changes in dielectric strength on the order of a few percent might occur, particularly at low concentrations of

chloride to be expected in these films (13). On the other hand,

increased film conductivity at certain active sites where grain

boundaries, inclusions, or other defects exist cause decrease in

local field strength and place the film at these sites under tension

with respect to the electrostricti ve compression forces (increased

as a result of ion adsorption) exerted elsewhere. Chloride ion,

or any of those species considered aggressive, could effect localized

film rupture after penetration to the metal surface by promoting

formation there of pockets of non-protective low density material

similar to that produced as a corrosion product during active pit

growth. As the pockets expand upon continuing oxidation at the metal/

oxide interface, further stress is placed on the film directly atop

the pockets, eventually resulting in fracture of the film. In iron,

for Instance, this low density material was identified as y-FeOOH (24).

This particular process of depassi vation , which we shall refer to as the spontaneous film rupture mechanism, provides a finite number of access sites for solution to contact unprotected metal. By suggesting

that alteration of film growth in the presence of sufficient concentration

of some detrimental species at the metal/film interface could increase

localized internal film stresses to a point where the tensile strength

of the film could be exceeded, we have modified the basic Sato mechanism

in such a way as to reasonably account for ion selectivity in

promotion of localized corrosion.

Susceptibility to localized corrosion will be determined by the

repassivation kinetics of the particular metal in the specific

environment which it contacts. Sustained metal dissolution, whether

resulting from slow film growth rates or from exceptionally high

anodic overvol tages , will lead to alterations in the environment

chemistry which are invariably detrimental to effective repassivation.

Thus, it becomes increasingly difficult to promote repair at the

local breakdown sites, and pitting corrosion can result.

If one accepts the hypothesis of dynamic breakdown and repair,

then all forms of localized corrosion can be considered as extremes

of one process. In stress corrosion cracking, applied tensile

stress not only increases the rate of breakdown events, but can increase the rate of dislocation movement and subsequent area of bare

metal production through slip step emergence. Rate of film repair

in that environment then determines rate of crack propagation.

Similarly, crevice corrosion results from the fact that mass transport

limitations over a rather extensive metal surface area result in

deleterious solution chemistry for repassivation. The process can rather quickly lead to sufficient solution acidity to attack the

remaining bulk surface film and cause depassivation on a large scale

-21-
with subsequent general attack beneath the crevice.

The fact that susceptibility to localized corrosion is related

to some critical limiting potential cannot be explained by a dynamic

breakdown and repair process unless those detrimental ions which

are involved in the repair process are either 1) specifically adsorbed

above this potential, thereby inhibiting film growth or 2) contribute

to the formation of non-protective films above critical potential

values

.

Repassi vation studies of iron in chloride solution by Ambrose

and Kruger (11) showed that film growth was normal below a critical

potential, but that films grown in the presence of chloride ion

above this potential were not protective.

The examination of our results on the effect of alloying on localized corrosion susceptibility of ferritic stainless steels

will be based on the following mechanism:

1) Aggressive ions (e.g., chloride) adsorb on and penetrate into

surface films at specific weak sites. Such sites occur at points of

higher film conductivity (around grain boundaries, inclusions, or other physical film defects) where the electric field is less, and

the film will be under tensile stresses. When normal film growth

is altered by the presence of chloride at the interface, growth

of the resulting low density oxidation product will eventually

cause rupture of the overlaying protective film.

2) The extent to which localized corrosion can proceed will depend

upon the repassi vation kinetics of the surface produced from film

rupture in the environment to which this bare metal will be exposed.

3) Sufficient deterioration of the environment due to metal dissolution (increased soluble corrosion products and lowered dH) will

at least affect film growth kinetics and in the extreme, inhibit it.

Pitting attack would result.

4) Application of external tensile stress will accomplish two

things: a) increase the number of breakdown sites on the metal

surface and b) increase the rate of slip step emergence leading to

a production of a number of sites with a relatively large surface

area of bare metal. Repassi vation kinetics will determine

susceptibility to stress corrosion cracking, provided rate of bare

metal production is of the appropriate value.

5) Creation of a crevice will provide a mass transport-limited

solution volume where soluble corrosion products and charge

balancing detrimental anion concentrations lead to a rapid increase

in film breakdown events and finally to a general dissolution of

the film within the crevice.

-22-
Effect of alloying on susceptibility to localized corrosion

Chromium additions

As Fig. 7 shows, an increase in chromium content in an alloy

increases the rate of repassi vation. Since chromium is known to

complex chloride (14) according to the reaction,

Cr(H20)^^ + CI" = CriH^O)^ Cl"^^ + H^O, [3]

it is conceivable that dissolution of chromium during depassivation

ties up chloride ions. The chloride ions are thereby prevented from

interferring with the rapid repassi vation of the metal surface.

What effect does Cr have on the depassivation process? Some

insight into this question can be gained from spectroscopic

el 1 ipsometri c studies by McBee and Kruger (15,16). They showed

that, when working at a few-specific wavelengths of light, additions

of chloride ion to the solution affected the optical parameters of

surface films on pure iron, and that such alterations which lead to

breakdown can be reversed b^ ,the chloride ion being removed from

the environment. When the;ch1or1de was removed, the optical

parameters returned to their previous values (Fig. 14). Such

results suggest that chloride ion is poorly bound in the film and

mobile enough to either move in or out depending on bulk solution

conditions.

When the same experiments were carried out using iron-chromium

alloys Instead of pure iron, a significant difference in behavior

was noted (16). First, much higher chloride concentrations were

necessary to cause breakdown, of the films on the Fe-Cr alloys than were used for the pure iron. Secondly, optical charges which occur

in the film on the Fe-Cr alloys during the period prior to breakdown

were not reversed when chloride ion was removed from the solution

as they were in the case of pure iron (Fig. 15). McBee and Kruger

attributed this difference to ability of the chloride ion to move via

vacancies in the crystalline films on pure iron, whereas chloride

ion had less mobility in the non-crystalline films on the Fe-Cr

alloys due to the greater possibility for covalent bonding in the

amorphous film. Once this film on the Fe-Cr alloy had been

contaminated with chloride ion, not only would it remain contaminated

even when the solution is changed, but if chloride containing solution

was reintroduced, McBee and Kruger found that breakdown occurred

Instantly.

The Cr additions thus operate in two ways to inhibit localized

corrosion. First, they retard alteration of the film's properties

that would lead to depassivation. As pointed out in the preceding

-23-
Fia. 14. Spectro-el 1 ipsometric measurements of changes in a and \i) for iron after introduction of chloride (t=0), and removal

of chloride from the solution, a, ii, and current are seen

to return to normal passive values once chloride ion is

removed. 410 mn wavelength light was used. From (15).
86 5

86 0

24

23

CBOO

>

OOqO o oo ° °

°

CHLORIDE

ADDED

~

Q O O O O

CMUOIieOE

RCMOVCO'

o o o

o o o o o o o o o o

o o

ooo

_ 00

J L

fe>20Cr

I L

0 20 40 60 to 100 ttO 140 t«0 100 200 220

Time , Minutes

Fig. 15. Spectro-ell ipsometric measurements of changes in A and

for Fe-20 Cr after introduction of chloride (t=0), and

removal of chloride from solution. A and ^ do not return

to passive values once chloride ions are removed. 450 nm wavelength light was used. From (15).
section, the most reasonable way in which the incorporation of

chloride in the film affects depassi vation is by changing the

mechanical properties of the film so that spontaneous film rupture

occurs. The second way the Cr affects localized corrosion, as shown

by the results depicted in Fig. 7, is in the promotion of rapid

repassi vation. It is only when a large supply of chloride ion

is present to retard repassi vation and promote depassi vation that

breakdown occurs. Chromium thus has a beneficial affect on repassi vation and in preventing depassivation.

Effect of Molybdenum

The role of molybdenum in increasing corrosion resistance of

ferritic stainless steels is an enigmatic one because there is some evidence that Mo is not present in the passive film (4). To

emphasize this point further, spectroscopic el 1 ipsometric experiments

on Fe-Cr-Mo alloys (20) show that molybdenum affects neither the

retention of chloride in the film nor does it alter the magnitude

of changes in ell ipsometric parameters A and i). Fig. 16 shows that

essentially the same change in optical properties occur when Mo is

present in an Fe-Cr alloy as those found when molybdenum is

absent (Fig. 15).^ These results, therefore, indicate that

molybdenum, unlike chromium, does not enhance the resistance of a passive film to undergo the breakdown that eventually leads to

depassivation.

On the other hand, the repassi vation process is accelerated

by alloying additions of molybdenum, particularly in the higher

chromium containing steels (see Figs. 2,3,4). Since molybdenum is

probably not present in the film, its effect on repassi vation must

depend upon its existence as a dissolved species in solution. The

experiments on the effect of potential on repassivation rate clearly

point to this possibility. As Fig. 8 shows when molybdenum is absent,

an increase in potential does not affect k, the rate of current

decay, a parameter related to the rate of repassivation. The current

increases with potential in this case and thus more alloy is dissolved

at the higher potentials. Even when molybdenum is present in the

alloy, the initial maximum current is increased with an increase

in potential and results in an increase in the concentration of

dissolved species (Fig. 9). Only in this case the presence of

dissolved molybdenum species affects the values of k, which increase

as the higher potentials provide increased molybdenum concentrations

in solution. We do not know what role the dissolved molybdenum

plays in repassivation. It is known to go into solution electrochemically in the +6 oxidation state (21,22) and to form molybdates

or polymolybdates . These are known to be corrosion inhibitors (19),

but how they enhance the rate of repassivation is unclear.

Spectroscopic ellipsometer measurements were the same for both

Fe-20 Cr and Fe-17 Cr alloys.

-26-
86 5

660

24

23

OCDOOCXD OOOOOOOOOOO o o o

CNUMIIOC

> AOOCO

s

t

00 oo o

qOOO-OOO o o o o

1 L

rfl7%Cf9%Mo

1 1 i 1 1 0 2040 «OMtOOtSOMOltOl00200 220

TiMC , MmuTES

Fig. 16. Spectro-ellipsometric measurements of changes in A and

for Fe-17 Cr-5 Mo after introduction of chloride (t=0),

and removal of chloride ions from solution. 450 nm wavelenath light was used. From (20).
Table II.

NaCI Solution Refractive L 4> 6A, c fU9 e; Concentration Index

i.v^^.o

l.ON 1.3432 113.86 29.38

2.5 1.3567 111.70 29.41 -2.16° +.03**
Whatever the origin of effect of molybdenum on repassivation

kinetics, the different rates of repassivation for alloys of

different molybdenum contents Iqad to different concentration

levels of soluble corrosion products in the crevice environment, and

the A changes measured during the inital phases of attack were

probably due to increases in refractive index of the solution. The

change in A is in the direction consistent with an increase in

concentration (Table II). Molybdenum, since it increases repassivation

rate, would decrease rate of change in refractive index. The point

in time at which refractive index, n, begins to range from its

initial value will obviously depend on chromium concentration in

the alloy, since this variable determines depassivation time. Indeed,

our results show that A (and solution refractive index) begins to

change at approximately the same time for fixed chromium concentration,

regardless of molybdenum concentration, but the rate of change is

proportional to molybdenum content for both Fe-14 Cr and Fe-17 Cr

(Figs. 11 and 12). Thus, the elapsed time when changes in A occur is a measure of the depassivation kinetics and rate of change of A

is a measure of the repassivation kinetics, the balance between

the two determining the nature and severity of the corrosion

attack on a metal surface.

CONCLUSIONS

1) An increase in the amount of molybdenum in Fe-Cr alloys

increases the rate of repassivation of bare surfaces in 1.0 N NaCl

solutions.

2) Molybdenum has no effect on the time to observe the beginning

of the processes leading to crevice corrosion, but the rate at

which these processes proceed is affected by the molybdenum content,

the more molybdenum the lower the rate.

3) If the onset of localized corrosion is assumed to be controlled

by a dynamic balance between depassivation and repassivation, the

molybdenum in ferritic stainless steels only affects repassivation

and not depassivation, while chromium affects both processes.

4) An addition of 1% Ni for an Fe-28 Cr-5 Mo alloy gives a 20%

reduction in its rate of repassivation in 1.0 N NaCl solution.

ACKNOWLEDGEMENTS

We are grateful to A. P. Bond and M. A. Streicher for providing

us with samples of alloys used in this study and to the Office of

Naval Research which supported this work under contract NAONR

18-89 NRO 36-082.

-29-
BIBLIOGRAPHY

(1) J. M. Kolotyrkin, Corrosion , 19, 261 t (1963).

(2) E. A. Lizlovs, "Localized Corrosion - Cause of Metal Failure",

ASTM Special Technical Publication 516, M. Henthrone, symposium

chairman, (ASTM, Tallahassee, 1972) p. 201.

(3) M. B. Rockel , Corrosion . 29, 393 (1973).

(4) J. B. Lumsden and R. W. Staehle, Scr. Met ., 6, 1205 (1972).

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

(6) J. R. Ambrose and J. Kruqer, J . Electrochem Soc . , 121 , 599

(1974).

(7) J. Kruqer and J. R. Ambrose, National Bureau of Standards

Report No. NBSIR 73-244, Technical Summary Report Number 4, Contract

NAONR 18-69, NR 036-082, 1973.

(8) M. Nagayama and M, Cohen, J. Electrochem. Soc , 109 , 781 (1962).

(9) M. A. Stretcher, Corrosion , 30, 77 (1974).

10) J. F. Bates, Corrosion , 29, 28 (1973).

11) J. R. Ambrose and J. Kruqer, Proc . Fourth Inti . Conq. on Met.

Corr . . N. E. Hamner, ed., Amsterdam. 1969 (NACE, Houston, 1972)

p. 698.

12) T. P. Hoar and W. R. Jacob, Nature , 216 , 1299 (1967).

13) H. S. Parker and C, A. Hardinq, J. Amer. Ceramic Soc , 53,

583 (1970). ' '

'

14) R. B. Heslop and P. L. 'Robinson, "Inorcanic Chemistry", (Elsevier,

Amsterdam, 1960) p. 444.

15) C. L. McBee and J. Kruger, Proc. IJ. R. Evans , Intl. Conq. on Localized Corr., Williamsburg, 1972 (NACE, in print).

16) J. Kruger and C. L. McBee, National Bureau of Standards Reoort

lC-90;. Contract OSW No. 14^30-2610, 1972.

17) G. Okamoto, Corr. Sci . , 471 (1973).

18) N. Sato, Electrochim. Acta , 16_, 1683 (1971 ).

19) I. L. Rozenfeld and L. V. Frolova, Dokl . Akad. Nauk. SSSR , 215,

1410 (1974).
(20) C. L. McBee and J. Kruger, unpublished work,

(21) J. k'. Johnson, C. H. Chi, C. K. Chen, and W. J. James,

Corrosion , 26^, 238 (1970).

(22) J. W. Johnson, M. S. Lee, and W. J. James, Corrosion , 26 , 507 (1970).

(23) P. Delahay, "Double Layer and Electrode Kinetics", (Interscience,

New York, 1965) p. 23.

(24) C. L. Foley, J. Kruqer, and C. J. Bechtoldt, J. Electrochem . Soc. , m, 994 (1967).
To be submitted for publication in Corrosion

Comnarison of Repassivation Measurements

to Constant Strain-Rate Tests in the

Stress Corrosion Cracking of 304 Stainless Steel

bv

J. R. Ambrose

Institute for Materials Research

National Bureau of Standards

Washinoton, D. C. 20234

ABSTRACT

Constant strain rate studies were performed usino AISI 304

stainless steel exposed to acidified l.ON NaCl solution (pH=3.0).

In this oarticular environment, reoassivation kinetics measurements

usina the tri bo-ell ipsometric technique had shown that the metal

dissolution rates were rather large due to reduced film repair

kinetics, indicatina the possibility of susceptibility to stress

corrosion crackinq (SCC). Reduced time to failure and maximum

sustained load at one strain rate suggest a relationship

between SCC, repassivation kinetics, and rate of pure metal

production.
To be submitted for publication in Corrosion

Comparison of Repassi vation Measurements

to Constant Strain-Rate Tests in the

Stress Corrosion Cracking of 304 Stainless Steel

by

J. R. Ambrose

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

Recent interest in the use of repassivation kinetics measurements to predict susceptibility of a material to stress corrosion

cracking in a particular environment has provided a large amount

of experimental data with which it is possible to estimate an extent of electrochemical reaction which occurs after the production

of bare metal (1,2). This laboratory has been involved in the

measurement of repassivation kinetics using the technique of

tri bo-el lipsometry (3,4). This technique is used to make optical

and electrochemical observations of repassivation of bare metal

produced by abrasion. It gives no information, however, about

the rate of bare metal production resulting from the film rupture

that occurs when a metal is stressed. Information obtained from

use of tri bo-el lipsometry about the repassivation ratio of metal

dissolution to film growth {R$)^ is, of course, essential in

determining the extent of reaction during one such incident and

thus giving a measure of susceptibility to SCC, but unless the rate

of bare metal production provides a more or less constant and

localized area of exposed metal, cracks cannot propagate. The use

of slow strain tensile machines which essentially measure the

effect of the rate of bare metal production on SCC susceptibility

has been described by Parkins (5) and already successfully used

by a number of workers (6,7). In this study, kinetics of repassivation

are compared to results obtained using the constant strain-rate

test. The rationale behind the use of the constant strain-rate test

is that if a particular material possesses a susceptibility to

stress corrosion cracking in a given environment, then there

should exist an optimum rate of bare metal production for propagation

of cracks to occur during straining. This would be determined by

the repassivation kinetics of that system. At strain rates much

MRf^) is the ratio of the total charge involved in repassivation to

that involved in protective film regrowth.

-33-
slower than this value, the bare metal produced will repassivate

without crack advance. Faster strain rates will reduce specimen cross section in a ductile manner before SCC can contribute.

Studies described here are in part based on repassi vation kinetics

measurements on AISI 304 stainless steel in 1.0 N NaCl solution (8).

One result of those studies was that the repassivation ratio, Rp, was sianificantly hiaher in an acidified 1.0 N NaCl (pH=3) at a potential

of +90 mV SHE (corrosion potential measured for 304 stainless steel

in this environment) than in neutral or basic solutions. This

svstem was therefore selected as a startinq point for the determination

of the relationships between repassivation kinetics, rate of bare

metal production, and stress corrosion crackinq susceptibility.

EXPERIMENTAL

Apparatus

Experimeni^ in this study were performed on the slow strain

tensile machine shown in Fig. 1 capable of producina strain rates

of 10-5 to 10-7 sec-1 . Load was measured potentiometrically with

a load cell

.

The electrochemical cell used in this study was fabricated from

teflon and utilized teflon Swagelok fittings to maintain liquid tight

seals (Fig. 2). The lubrication properties of teflon prevent

transmission of a portion of the applied load to the cell itself

during straining. The volume of the cell was about 50 cm^. The

auxiliary electrode was housed in a separate cell to nrevent build

up of counter electrode reaction products.

Material

Specimens were machined from 17.78 cm lengths of 6.35 mm diameter

cold drawn annealed AISI type 304 stainless steel rod stock and had

reduced cross sections with 2.54 cm radius shoulders which gave a 2.54 cm long guage length 3.175 mm in diameter. Before placement in the

electrochemical cell, specimens were degreased with spectro-grade

methanol, rinsed in distilled water and dried in a jet of compressed

air.

Experiments were conducted in 1.0 N NaCl solution acidified to

pH 3 with 1.0 N HCl . Solutions were oreoared usi no reagent arade

chemicals and distilled water (7.0 x 10-7 ohm-' cm"'). Where load or

time to failure data are reported, values are normalized with respect

to air. All experiments were performed in duplicate and recal ibrations

I am grateful to R. N. Parkins and B. E. Wilde for their contribution

to the design of this instrument.
Fig. 1. Photoqraph of the constant strain rate machine used in these

studies.

-35-
Fig. 2. Photograph of the teflon electrochemical cell (unassembled)

used in these studies.

-36-
(tensile tests in air at a given strain rate) were repeated when time

to failure or sustained loads varied by more than 1% from the air

fracture value.

RESULTS AND DISCUSSION

Changes in specimen potential with time were measured during

straining at 2.2 x 10"^ sec"^ for 304 stainless steel in air

saturated 1.0 N NaCl solution (pH 3.0) and results are shown in Fig.

3. The initial potential increase is probably due to depletion

of oxygen within the sealed cell. The final potential jump occurred

at specimen fracture.

Stress corrosion susceptibility tests were then run in the

acidified NaCl solution at several different strain rates using

specimens polarized to. +90 mV SHE. This environment was chosen

because previous work (8) showed that the repassivation kinetics

observed appeared to be favorable for SCC susceptibility. It can

be seen from Fig. 4 that the repassivation ratio (Rp) determined

by tribo-ell ipsometry was of sufficient magnitude (at pH 3, +90

mV SHE) to indicate these conditions as the more likely of those

studied for observing susceptibility to SCC.

Maximum load and-^-time to failure data, normalized with

respect to air, are given in Fig. 5 and in Table I.. A scanning

electron microscope was used to examine the fracture surfaces

of both air fracture and in situ fractures at the maximum strain

rate obtained with the equipment oresently available. From Table I,

it can be seen that the reduction in maximum load and time to failure

indicate possibility of SCC attack. The fracture surface of this

specimen was examined with a scanning electron microscope. The

central portion of the fracture surface shown in Fin. 6 is of the

dimpled nature characteristic of ductile tensile failure. However,

one area of the surface (see arrow. Fig. 6) showed a markedly

different topoqraphy (Fig. 7) which, under higher magnification,

shows some evidence of the cleavage characteristic of stress corrosion cracking in austenitic stainless steels (Fia. 8). The fracture

surface is pock-marked with two distinctly different types of holes —one similar to that produced during ductile fracture, and the other,

crystallographic in appearance, probably the initial stages in

formation of pits on the newly generated metal surfaces (Fig. 9).

Slow strain exoeriments v/ere performed at the same potential

and strain rate on 304 stainless steel specimens exposed to neutral

1.0 N NaNO^cn^ environment in which the reoassivation ratio was sianificantly lower. No change in maximum apolied load or time to

failure with respect to air was detected.
Fiq. 3. Changes in potential (mV SCE) and sustained load versus time

durinq constant strain rate experiment [2.2 x lO'*^ sec"l] for

304 stainless steel in acidified 1 .ON MaCl solution (pH = 3.0).

-36-
10 Repassivation Ratio

Effect of pH

pH=3

mcoul

nm

pH=ll

10

• — ' 0 10 lOC

msec

Fiq 4 [From (8)1. Chahaes in Repassivation Ratio, Rp (Ioqiq Rp vs loam ms) for 304 stainless steel in air saturated l.ON NaCl

(pH 3, 7, and 11) at +090 mV (SHE) followina removal of

surface film by abrasion at -1000 mV (SHE).
Table I.

strain rate, sec"^ time to failure, hrs. max. sustained load, kq.

l.ON NaCl (pH=3) air 1 .ON NaCl (pH=3) air

2.1 X 10-^ 870 875 2950 . 2955

2.1 X 10'^ 75.1 75 .0 3000 3000

6.2 X 10"^ 35.0 34 .5 2960 3080

8.1 X 10-^ 24.5 24 .5 2900 2900

8.3 X 10'^ 21.0 22 .2 2900 3260
1.0 - aa>3

O

Li.

0)E

c

ouo

0)

>

*—o

a:

0 75 O—O REL. RED IN TIME TO FAILURE

a-- -A REL. LOAD RED.

0 50

10 -6

Strain Rate, sec"'

JO

<

COo

D

0.75 Q3]

fDQ.

Co

5'3

10 0 5C

5

Fiq. 5. Chanaes in time to failure and maximum sustained load (both

normalized with respect to air) versus strain race for 304

stainless steel in acidified 1 .ON NaCl solution (pH = 3.0).
Fiq. 6. Scanning electron micrograph (x20) of the fracture surface of

a 304 stainless steel soecimen fractured at a constant strain

rate of 8.3 x 10'° sec-l in acidified 1 .ON NaCl solution (pH

= 3.0),
Fig. 7. Scannina electron microqraoh (x500) of the area of the fracture

surface indicated by arrow in Fiq. 6. Specimen is 304 stainless

steel fractured at a constant strain rate of 8.3 x 10"^ sec"^

in acidified l.ON NaCl solution (pH = 3.0).
Fia. 8. Scanninq electron microqraph (x2000) showing areas resembling

cleavage on fracture surface of 304 stainless steel fractured

at a constant strain rate of 8.6 x 10"" sec"^ in acidified

l.ON NaCl solution (pH = 3.0).
Fiq, 9. Scannino electron micrograoh (x5000) showing the types

of Dits produced on the fracture surface of 304 stainless

steel fractured at a constant strain rate of 8.6 x 10"^

sec"' in acidified 1 .ON NaCl solution (pH = 3.0).
These results indicate that the effect of potentiostatic exposure

of 304 stainless steel in acidified sodium chloride solution on time

to failure and maximum sustained load during slow straining must arise

from a metal removal component occurring in conjunction with ductile

tensile failure during the time span of the experiment. Anodic current

density measurements made during the duration of exposure lead to a calculation of an average depth of metal penetration of 5 x 10-6 cm which cannot account for the measured decrease in maximum sustained

load. Of course, it is conceivable that localized attack in the

form of pitting could reduce the effective cross sectional specimen

area to such an extent as to reduce the sustained load. While

examples of pitting were observed in the SEM micrographs, the

crystal lographic pit density was low. The evidence of some cleavage

in the same SEM micrographs lends better support to the conclusion

that a sec process is operative under these conditions.

Although definitive conclusions must await further experimentation,

it is interesting to note that on the basis of relative repass ivati on kinetics of 304 stainless steel in neutral and acidified sodium

chloride solution, a greater susceptibility to SCC seemed likely

in the acid solution because this environment produced lower film

growth kinetics. Although the still limited evidence reported

here does not confirm a marked tendency for cracking in acidified

chloride environments at any finite propagation velocity under

service, it does suggest replacement of conventional susceptible/

non susceptible criteria which, while valid for predicting failures

that occur in short time spans, are not suitable for measuring

susceptibility when crack propagation is very slow. For practical

situations where failures take years to occur such slow crack

growth rates are still important. Here, especially for ductile

alloys where fracture mechanics techniques present problems, the

combination of repassi vation kinetics measurements and constant

strain-rate tests may be a valuable failure predictive aid.

Many questions remain, however. For example:

1) What is the effect of using higher strain rates than are presently possible with this apparatus? According to Parkins (5)

it is to be expected that relative time to failure and sustained

load values will reach a minimum before returning to unity at sufficiently rapid strain rates where material fails ductilely before

SCC can be expected to occur.

2) What is the effect of potential on strain experiment

parameters? One should find a point at which failure by general

anodic dissolution or pitting will take over.

3) From Fig. 3 [Ref. (8)] it can be seen that the repassivation

ratio in a pH 11 solution, although less than for the pH 3, is

somewhat greater than for neutral solutions. What should be the

effect of exposure to this environment on the tensile properties

during slow strain?

-46-
The work, is presently being directed along these lines.

ACKNOWLEDGEMENT

I gratefully acknowledge the support of the Office of Naval

Research who sponsored this work under Contract No. NAONR 18-69, NR

035-082.
BIBLIOGRAPHY

R. B. Dieqle and D. A. Vennilyea, J . Electrochem Soc . , in print.

T. R. Beck, Electrochimica Acta , 18, 815 (1973).

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

J. R. Ambrose and J. Kruger, J . Electrochem. Soc . , 121 , 599 (1974)

R. N. Parkins, "The Theory of Stress Corrosion Cracking in Alloys"

J. C. Scully, ed., (NATO, Brussels, 1971) p. 449.

J. M. Sutcliffe, R. R. Fessler, W. K. Boyd, and R. N. Parkins,

Corrosion , 28, 313 (1972).

D. C. Deeqan and B. E. Wilde, Corrosion , 29, 310 (1973).

J. Kruaer and J. R. Ambrose, National Bureau of Standards Report

No. NBSIR 73-244, Technical Report No. 4, Contract NAONR 18-69,

NR 036-082.
To be published in the Proceedinqs of the

20th Meeting of the Materials Failure Prevention Group

Failure by Stress Corrosion Crackinq - Current Approaches

Toward Failure Prediction

by

J . Kruqer

National Bureau of Standards

Washinoton, D. C. 20234

ABSTRACT

Stress corrosion crackinq (SCC) produces failures in a material

when it is subjected to the combined effects of mechanical stress

and reaction with an environment. All proposed mechanisms of SCC

seek to explain how the combination of a given level of stress, a particular material, and a qiven environment can lead to the

initiation and oropagation of cracks. The three major cateoories

of mechanisms that are qenerally proposed are (1) active path

dissolution, (2) stress-sorption, and (3) embrittlement.

In recent vears attempts to determine the mechanisms have

been enhanced by new concepts and techniques that flow from the

three major scientific disciplines that underlie the complex

problem of SCC - chemistry, metallurgy, and mechanics. From

chemistry have come two main currents: (a) increased awareness

of the importance of characterizing the altered environment inside

a arowing crack, and (b) a recognition of the importance of the

regrowth rate of a protective film on a bare surface exnosed

when that film is broken by stress. The main emphasis in the

metallurgy of SCC has been on the crucial role played by structure

in general, and particularly at the tip of a crack. Aiding this

objective has been the use of the high voltage electron microscope

which can directly look at the interaction of the environment with

structural defects, e.g., dislocations. Finally, the major thrust

in the mechanics of SCC has been the application of the concepts

and techniques of fracture mechanics.

The impact of these new concepts and measurement techniques

on predictinq and preventinq SCC failure will be discussed.
To be published in the, Proceedings of the

20th Meetinn of the Materials' Failure Prevention Group

Failure by Stress Corrosion Cracking - Current Approaches

Toward Failure Prediction

by

J. Kruger

National Bureau of Standards

Washington, D. C. 20234

Far too often the failure of a metal part involves an alloy that was

chosen because of its good corrosion resistance to the environment in

which it is designed to operate. Moreover, it failed at stress levels

well below its normal fracture stress. More than likely, this and

similar such unexpected failures were due to stress corrosion cracking

(sec). The insidious nature of this mode of failure results from the

situation just described where, in the absence of the possibility for

sec, one would expect both good fracture and corrosion resistance. Fig.

1 from a paper by Staehle (1) illustrates how SCC markedly reduces the

load carrying cross section of a part over that to be expected from

environmental attack (corrosion) alone or from purely mechanical failure.

Thus, it is from this synergistic interaction of corrosion and stress

that SCC arises, and the definition of the phenomenon reflects this fact.

For example, the ASTM Committee on the Corrosion of Metals defines it as

"a cracking process requiring the simultaneous action of a corrodent and

sustained tensile stress. This excludes corrosion-reduced sections which

fail by fast fracture. It also excludes intercrystalline or transcrystal

line corrosion which can disintegrate an alloy without either applied or residual stress." This definition embraces many of the other terms that

focus on particular mechanisms such as "hydrogen embrittlement," "caustic

embri ttlement," "season cracking," etc.* This diversity of terms has

tended to produce confusion. Many workers, for example Staehle (1) and

Brown (2), have proposed that the term "stress corrosion cracking" be

used as a generic term to avoid entangling and confusing the macro

phenomenon with any special mechanism that may govern a particular

situation. This proposal is being increasingly adopted in the modern

literature.

Because of the diversity of mechanisms affecting SCC, and because

it takes place in what are thought to be innocuous environments, this

mode of failure was not generally considered too important as late as 1940.

*It does not include "corrosion fatigue" which involves a cyclic stress

rather than a sustained static stress.
oiide or rust

Fia. 1. A comparison in the reductions in load-carrying cross section

as a result of (a) general corrosion and (b) stress corrosion.

The effect of stress corrosion on the tensile curve is shown

in (c). From Staehle (1).
In a discussion accompanying a 1940 paper by Hodge and Miller (3) which

clearly highlighted the problems of SCC, Brooks pointed out that "... the

industrial implications of this phenomenon do not loom large. Failures

are rare and will become more so as understanding spreads " Not only

were his predictions not realized, but the systems where SCC is recognized

as the cause of a failure have increased markedly throughout this century.

It was first recognized at the end of the 19th century when brass

cartridge cases (containing residual stress) failed because of exposure

to the ammonia present in stables. Table I documents the marked increase

in the recognition of SCC as a source of failure, going from the days of

transportation by horse-drawn vehicles to the space vehicles of today.

The number of systems where previously failures were not expected continues

to grow, especially as greater demands are made on materials.

Because of the increased identification of SCC as a major cause of

failures, there has resulted a significant increase in research directed

towards understanding the phenomenon and developing mechanisms to explain

it. Two recent proceedings of major meetings (5,6) and a forthcoming one

(7) give excellent accounts of the current state of thinking on the origins

of SCC. This paper will briefly outline the major mechanisms that are currently believed to explain SCC and then show how these mechanisms have

pointed the way to the development of new experimental approaches which,

besides enhancing our understanding of SCC, have,more importantly, provided us with opportunities for the development of new capacities for

failure prediction and prevention.

Mechanisms

In order to consider the important mechanisms that have been proposed

to explain the origins of SCC, we must first list the characteristics,

most of which must be present to produce this mode of failure. This is

necessary because a valid mechanism must explain the necessity for these

characteristics to be present if SCC is to occur. Brown (4) has listed

the following features generally common to SCC:

1) The existence of a tensile stress either applied or present as a residual stress.

2) The alloy usually shows good resistance to corrosive attack by the

environment causing cracking. This is why alloys chosen because of their

corrosion resistance to an environment may be subject to SCC.

3) SCC for a given alloy usually occurs more readily in a rather specific

environment. Moreover, the key damaging species in that environment need

not be present in large concentrations. Recent studies (9,10) have indicated

that the dependence of SCC on environmental specificity is not as pronounced

as originally thought in early SCC studies.

4) The microscopic appearance of stress corrosion cracks is that of a brittle fracture in spite of the fact that the metal itself may still be

quite ductile in a mechanical test. The cracks frequently exhibit multiple

branching, but this does not always occur.
Table I

WHEN

HISTORY OF STRESS CORROSION PROBLEMS

[Adapted from Brown (4)]

WHAT

Late 19th Century

End 19th Century

Early 20th Century

Early 20th Century

1930's

1930's

1950's

1 950-60 's

Brass in ammonia

Boiler steels in caustics

Mild steels in nitrates

Aluminum alloys in moist atmospheres

Stainless steels in chlorides

Magnesium alloys in moist atmospheres

Martensitic steels in aerospace environ.

Titanium - Hot salts, N20^, methanol
5) In order for SCC to occur a threshold stress or stress intensity

must be exceeded.

To Brown's list must be added another characteristic becoming

increasingly recognized (9,11). There exists a critical potential. Any

metal in an environment where its potential lies below this critical

potential is thought by some to be immune from SCC. This is a more recently proposed characteristic that has been strongly advocated by

Uhlig (11).

At one time it was believed that SCC could not occur for pure metals,

but some recent work (13,14) has found SCC for high purity copper. As

with all the other characteristics, one can usually find exceptions.

Using these characteristics of SCC as a basis, a number of mechanisms

have been developed. At one time some believed that there existed a universal mechanism that applied to all systems where SCC was found. It

is now widely held that SCC can be due to a number of different mechanisms,

and that there are alloy-environment systems where one or more of the

mechanisms outlined here may apply. Of the many proposed, three major

mechanisms are currently the ones most actively used. Moreover, the

others proposed can be considered as variations or sub-sets of the three

major ones. All of the mechanisms involve the complex interactions of

the three disciplines that control SCC, chemistry, (electrochemistry),

metallurgy, and mechanics. None of the mechanisms will be discussed in

great detail, and a critical evaluation of their pros and cons is outside

the scope of this paper. Instead these mechanisms will be briefly described, their strong and weak points detailed, their failure prevention

implications indicated, and, in the next section of this paper, the

techniques that they point to as possible tools for failure prediction

will be discussed.

Stress-Sorption Mechanism - This mechanism, most actively advocated

by Uhlig (14,15), proposes that a specific species adsorbs and interacts

with strained bonds at the crack tip causing a reduction in bond strength,

or, thermodynamical ly stated, a reduction in the surface energy, y> which

leads to a lowering in the stress required to produce brittle fracture.

The relationship between the theoretical stress to produce brittle fracture,

a^, and Y is given by the well known expression

Of = [1]

where E is the elastic modulus and d is the atomic spacing. The adsorption

of a damaging species at a crack tip lowers af and therefore the threshold

stress for SCC is the minimum value to which y is lowered by adsorption.

Fig. 2 outlines the mechanism schematically. Uhlig proposes that an important factor arguing for this mechanism is the existence of a critical

potential below which SCC does not occur. He explains that the effect of

potential can be attributed to the strong dependence of adsorption on potential. He bases this on many studies by electrochemists who have found
(a) Sub -> CRITICAL (b) Adsorption Reduces

Stress Bond Strength

Fig. 2. Stress sorption mechanism for SCC.

(a) chemisorption of ion at tip of crack.

(b) weakening of interatomic bond at the tip of the crack

so that applied tensile stress leads to the breaking

of this bond.
that the adsorption of an anion (usually the damaging species for SCC)

takes place at potentials above the point of zero charge for a given

surface in a given environment. (This strong dependence of failure or potential explains why SCC can be prevented by cathodic protection in some systems). Environment plays a role in the stress-sorption mechanism

because of the specificity of adsorption of a particular species on a surface, and because environment brings the potential to a value where

adsorption can or cannot take place. Moreover, if one can introduce

into the environment an extraneous species that displaces the damaging

bond-weakening adsorbed species, cracking becomes more difficult. Uhlig

(14) gives examples where such extraneous species inhibit SCC. The stresssorption mechanism is more widely accepted in the case of liquid metal

embrittlement or stress cracking of plastics. In the former case, for

example, Preece (15) points out that the theory is successful in explaining

the observations of a brittle to ductile transition with increase in

temperature (desorption) , an "instantaneous" failure mode, the increase

in severity of embrittlement with increase in strain rate and flow strength,

and, possibly, specificity.

For SCC in aqueous ^efivironments , however, a number of problems with

the stress sorption mechanism have been pointed out (9,10). Among them

are the following:

1) In the cracking of ductile ci^lloys, the ductility present is sufficient

to produce extensive plastic deformation at a crack tip. The blunting that

this produces eliminates the infinite sharpness required for this mechanism

to work.

2) The large amount of metal dissolution observed in many straining

electrode experiments (to be discussed later) poses the question of what

happens when the metal dissolves out from under the adsorbed film.

3) Crack growth rates are considerably slower than one would expect

from this model

.

4) Studies (17) have shown that thin protective films exist on metals

in SCC environments, and that these films reform rapidly, on the order of

milliseconds in some cases. The weakening of metal bonds require that

rapid direct access to the base metal surface is necessary, but the

existence of these films would prevent or strongly influence this

adsorption process.

The measures suggested by the stress-sorption mechanism to promote

SCC failure avoidance are as follows:

1) Lower the potential of a metal to be protected below the critical

potential for SCC either by coupling another more active metal to it or

by cathodic protection using an applied current.

2) Introduce inhibiting anions in the environment.
To properly apply these measures requires techniques that enable

both the determination of the critical potential for SCC and the potential

of the metal in an environment. The latter requirement is not as easily

met since the environment that counts is that environment existing in

flaws or cracks. In addition to a realistic characterization of the

environment, there exists an equal need for realistic characterization

of the stress system acting on a crack tip, since the model requires

adsorption at strained bonds or dislocations. Thus, techniques that yield

both the geometry of stress application and expected dislocation

configuration would be valuable for failure prediction using the guidelines

provided by this mechanism. Finally, ways to measure the rate of adsorption

of the relevant species as a function of stress geometry and potential

would be invaluable if this mechanism has merit.

Film Rupture-Metal Dissolution Mechanism - This mechanism was originally suggested by Logan (18) and Champion (19). It has also been

refined and extended by a number of workers (10,17,20-24). The mechanism

proposes that the protective film normally present on a metal surface

(if it were not present the metal would fail by corrosion alone) is

ruptured by continued plastic deformation at a crack tip where the

exposed metal becomes a very small and restricted anodic region where

metal dissolution takes place. The rest of the metal surface, especially

the walls of the crack acts as a cathode. Susceptibility depends on the

rate at which the metal exposed by film rupture is repassivated vis a v1s

the rate of metal dissolution. Fig. 3 gives a schematic representation

of this mechanism. A clear picture of this process has been provided by

Vermilyea (24) in a recent paper where he lists the following five steps

in the process that leads to crack extension:

1) The film is ruptured at time zero.

2) Crack growth by dissolution occurs for a distance of j- during the

time of repassi vation t^.

3) The dissolution process removes the material at the old crack tip

and thereby creates a strain transient just ahead of the newly created

crack tip as shown in Fig. 4.

4) At a time tL a critical strain builds up to a value sufficient to cause rupture of the newly reformed film at the base of the new crack.

j|

5) The process repeats itself, the crack propagation rate, ^5 given by

Vermilyea as
Stress _ , Stress _ _ Stress

(a) FILM (b) METAL (c) REPASSIVATION

RUPTURE DISSOLUTION

Fia. 3. Film rupture-metal dissolution mechanism for SCC.

(a) RuDture of the protective film at tip of a crack by the

application of a tensile stress.

(b) Dissolution of the metal exposed when film is ruptured.

Tip of crack is an anode and film covered walls of the

crack is a cathode.

(c) Repassivation of the exposed metal at the tip of the crack.
STRAIN

Fig. 4. Time deoendence of events occurring after film rupture.

Repassivation time is t^. The time to produce the critical

strain, e^, necessary to cause film rupture is t^,. From

Vermilyea (24).
A sub-set or variation of the active path dissolution mechanism just

described is the brittle-film mechanism (25,26). It differs from the

dissolution mechanism in that crack advance is mainly by the production

of a penetrating brittle film which undergoes fracture when tensile

stress is applied. The crack produced enters the ductile metal below

this brittle film and a new layer is formed on the bare metal so produced.

The process continues by a repetition of the film fracture-reformation

cycle.

One can conceive of at least three factors that affect the

general mechanism.

1) The rate of production of bare sites, which is dependent on the

magnitude of the stresses applied, the metallurgy of the material, and

the geometry of the stressed system tthe value of t^, is important here).

2) The rate at which the bare sites produced by the applied stress are repassivated (t|_ is important here).

3) The resistance of the protective film (its ductility) to rupture

[see (24)], (again t^ is important).

The nature of the species in the environment has an effect on tp, or probably more importantly, the rate of repassivation (17). This is so because small impurities can affect the composition of the newly reforming

film and rates of growth are sensitive to the presence of impurities (27).

Also, environment can affect tc through its influence on film ductility

(28). Potential affects both film growth kinetics and metal dissolution

rate so that the balance of these two competing factors is important, the

latter affecting L, the distance of crack advance. The concept of a critical potential for SCC is as valid for the film rupture-metal dissolution

model as it is for the stress-sorption mechanism.

There are a number of facts that create problems in accepting the

film rupture-metal dissolution model for all cases of SCC. Among the

most important are the following:

1) The quasi -cleavage character of some SCC transgranular fracture

surfaces indicate brittle failure rather than dissolution.

» .•'./ ' 2) Crack propagation occurs at a surface which is itself rapidly

dissolving at a high rate (29).

3) There is a possible discrepancy between crack growth rate and the

currents observed in straining electrode experiments which purport to measure the current involved in dissolution at crack tips (22).

The measures suggested by the active path dissolution model to promote

SCC failure avoidance will include one suggested for the stress sorption

model, i.e., control- of potential to a value below the critical potential

for SCC. Another measure is the control of repassivation kinetics by

environment or alloy modification. In order to apply the latter measure

for SCC avoidance, techniques for measuring repassivation kinetics as a function of potential and environment are required. However, in order to make

such measurements one must, as suggested earlier for stress-sorption,

-60-
be able to characterize the environment and potential existing in cracks.

Finally, and most difficult, one must have a technique for determining

Iq, the time required to huild up a rupturing stress in a film as a function of the metallurgy of the alloy whose failure we wish to prevent.

This requires the measurement of the ductility of the protective film

on an alloy and the geometry of the stress system acting on a crack.

Hydrogen-Embri ttlement Mechanism - This is considered by some to be

a different phenomenon from SCC, but since cathodic reactions (which

produce hydrogen under the proper conditions) are as much a part of

corrosion as anodic ones (producing films and/or metal dissolution), it

is difficult to exclude it in a discussion of SCC. In this mechanism

fracture results from the production of a brittle region at the crack

tip because of the introduction of hydrogen into the alloy via cathodic

reactions (see Fig. 5). The production of this brittle region has been

attributed to a number of causes among which decohesion (30), new phase

formation (31-33), and pinning of dislocations (34) can be mentioned.

The environment can affect the process by influencing the rate of production of hydrogen ions and by enhancing or inhibiting the entry of

hydrogen into the metal by the interposition of adsorbed or reaction layer

films. In aqueous systems the potential at the crack tip determines

whether the production of hydrogen at the surface is thermodynamical ly

feasible. The composition of the environment influences what potential

is necessary to produce hydrogen, its rate of production, and by means

of the presence of inhibiting or enhancing species, whether it can readily

enter the metal lattice. In gaseous systems, hydrogen pressure plays the same role as potential in aqueous environments. Inhibiting species, e.g.,

oxygen, can also be added to gases (35).

It has been suggested by some workers, for example, Troiano and cov^orkers (36), that almost all aqueous SCC is hydrogen embri ttlement.

They found that austenitic stainless steels, which have been assumed to

be non-susceptible to hydrogen embri ttlement , could be made less ductile

by hydrogen charging, and that one could observe the permeation of hydrogen

at positive currents (where one should only expect metal dissolution).

Other workers (31-33,37-39) have found the formation of a phase or phases

capable of undergoing hydrogen embri ttlement.

All of these findings do point to greater universality of application

for the hydrogen embri ttlement mechanism. However, as always, there are counter arguments that rule against such universality. Recent v/ork by

Wilde and Kim (40) showed that for an austenitic stainless steel while

rate of hydrogen permeation was higher at negative potentials, time to

failure was also higher at these hydrogen producing potentials. This

finding also agrees with results that show that cathodic protection

(which should promote the production of hydrogen) does retard SCC for a number of alloys (16). Indeed, to confuse the picture further, it has

been found by Bernstein and Pickering (41) that anodic dissolution can occur in the cracks of cathodically polarized ferrous alloys.
Stress

Cathodic

Reaction

H"^+e-^ H

^ M++e

Anodic

Reaction

Production of

Brittle Region

OR Path

Fig. 5. Hydroqen embri ttl ement mechanism for SCC. Cathodic reduction

of H+ at the tip of a crack produces hydroqen which enters

the vicinity of the tip and embrittles the metal immediately

ahead of an advancing crack.

-62-
This difficulty in determining whether hydrogen embrittlement is an operative mechanism is most important in developing measures for failure

prevention. For example, unless one knows the operative mechanism, a change of the potential to what is believed to be a protective value may

actually accelerate failure. Obviously techniques that determine what

are the' real potential and environment at the tip of the crack are crucial.

Likewise, electron micrography of fracture silVfaces which help in

determining the mode of failure and can thereby suggest the proper prevention

measures are also of utmost importance.

Techniques Suggested by Mechanisms

In the preceeding section the discussion of the major mechanisms

of sec pointed to the types of information that are needed in order to

develop our -ability to predict SCC susceptibility and to institute

suitable measures for prevention. In this section the techniques

developed in recent years, in many instances under the influence and

stimulation of the concepts of these mechanisms, will be discussed.

The techniques brouqht into being in tnis manner have their roots in

the three broad ^disciplines that control the phenomenon of SCC: chemistry,

metallurgy, and mechanics.

Techniques Dealing with the Cheiin'scry of SCC - Two areas have been

attacked with increasing interest in Vecsnl years . As pointed out in

the previous section they are both important in all three major mechanisms

although not always well recognized. Ihey are the techniques directed

towards determining the environment in cracks arid the kinetics of repassivation.

a) The environment in cracks - While it has been recognized for a number of years that the environment 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, brown and co-workers (42-44), using the insights

provided by the large body of work of Pourbaix [see the review in (45)

J

that systematized the role cf pH and potential in all corrosion processes,

set about to measure these parameters in SCC cracks. Their original

approach was to immobilize the corrodent within SCC cracks by freezing

it in liquid nitrogen, breaking open the crack mechanically, and then

analyzijig the thawed out solution removed from the crack. A later

development is illustrated in Fig. 6 where microelectrodes 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 (46) and the combination of a capillary technique with thin

layer chromatography (47).

The value of these crack cherhistry techniques becomes evident

when Fig. 7 is inspected. It shows for a large number of steels of

different compositions in environments of either hign or low pH th-at

the potential and pH inside the crack falls in a region of the pll-potential diagram (Pourbaix diaaram) where one expects both corrosion and,

./ 3-
Fig. 6. Experimental arrangement for determining the acidity (pH)

and potential at a crack tip. The glass and reference

electrodes measure the pH electrochemically by sampling

the solution that wets the lens tissue after the solution

travels the length of the crack. From Smith, Peterson,

and Brown (44)

.
Fiq. 7. Potential -pH diaqram showing the potentiaUpH values in

the crack tip environments for a number of high strength

steels. From Brown in Ref. (6), p. 197.
because the points on Fig. 7 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-realistic conditions and lead to an improper choice. A good discussion of the electrochemical ramifications

of knowing the proper environments in a crack is given by Pourbaix (45).

In addition to knowing the nature of the solution in a crack and the potential

at the crack tip, it is of great importance to know another aspect of

the events occurring at the crack tip, e.g., how much hydrogen is entering the

metal lattice? As the discussion on the hydrogien embrittlement mechanism

stresses, this knowledge is of crucial importance. A tool for obtaining

this information has been developed based on the work of Devanathan and

Stachurski (48). By using a thin membrane of the metal . studied it is

possible to measure the amount of hydrogen that is evolved on one side of

the membrane when it passes through the metal by oxidizing it electrochemical ly in a chamber on the other side of the membrane. The anodic

current density at steady state involved in the oxidation of the permeated

hydrogen is a measure of the amount of hydrogen entering the lattice on

the hydrogen evolution side of the membrane. A discussion of the technique,

its difficulties, and validity has been qiven recently by Bockris (49).

Fig. 8 from a paper by Kim and Wilde (40) illustrates the importance

of this technique. It shows that the rate of hydrogen penetration cannot

always be related to SCC susceptibility.

b) Repass ivati on Kinetics - In all the mechanisms proposed for SCC,

access of the environment must be provided to the metal surface. This

surface, if it is suitable for withstanding ordinary corrosion, usually

has a protective film. Therefore, as discussed earlier, the breaking

of this film and -the rate at which it regrows (repassivation kinetics)

are of great imoortance in determining the time during which the environment

can interact with the bare metal surface. in any one of the different ways

the various mechanisms propose.

Scully (20,50) 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 metal -dissolution mechanism

because coupled 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 prooaqate. 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 (10) 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 (51-56). 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

-66-
HYDROGEN ABSORBED (mC/5cm^)

20

TO

Type 304 Stainless Steel Temp 146 C

• Sat'd Lithium Chloride

+1 g/l Na3As03

3 -4 5 6

I CATHOOIC

6 5 4

ANODIC

LOGARITHM APPLIED CURRENT DENSITY (A/cm^)

Fig. 8. The amount of hydroqen absorbed in 304 stainless steel as

a function of applied current. Na3AsO-3 (curve B) enables

qreater amounts of hydroqen to enter the metal. In spite

of the fact that more hvdrogen enters the metal in the

cathodic region no SCC failures were observed for currents

qreater than a cathodic current of around 10"^ A/cm^. From

Wilde and Kim (40).
the dissolution-repassivation event that occurs subsequent to exposure

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

repassivation (Fig. 9). In these mainly electrochemical techniques, one cannot separate out the current involved in dissolution from that

involved in passive film formation. A technique recently developed by

Ambrose and Kruger (56), however, does attempt to measure separately the

film regrowth transient by optical means (ellipsometry) and to compare

it to the overall dissolution-repassivation current. With this technique,

called tribo-ellipsometry 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 current as shown in Fig. 10. In Fig. 10 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 non- susceptible nitrate solution. With tribo-ellipsometry one can determine

the ration of the total current to that involved in repassivation only.

This ratio Rp* when the exposed surface area is known is given by the

expression

where Qd = charge involved in dissolution and Qy. = charge involved in

repassivation.

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

• Work on the cracking of low carbon steel (17) has shown that SCC does not occur when the effectiveness of repassivation is neither too high nor too

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

film Rp* is 2.8 in 1 N Na NO2 (25°C) where it is nonsusceptible; Rp* is

26 in 4 N HaHO^ (90°C) where it does undergo SCC, and Rp* 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.

This section, devoted to chemistry, has described techniques for

evaluating the role of the environmental side of the equation; the next

two sections are concerned with the other side of the equation.

Techniques Dealing with the Metallurgy of SCC - In recent years two

experimental approaches have become increasingly important in providing

insights into the role of metallurgy in SCC. They are constant strain

rate tests and a variety of applications of electron microscopy.

a) Constant Strain Rate Tests - This test, which produces experimental

data as depicted in Fig. 11, can be applied both to mechanistic studies

as well as used for a routine testinq. The technique has been applied

to studies of mild steels (59), Ti alloys (60), and Mg alloys (23). The

mechanistic importance of the technique is that it ties the metal lurgical ly

-68-
A comparison of oscilloscope repassivation transients for

Ti-8 Al-1 Mo-1 V alloy in IN NaCl and IN NaNOo solutions at

+706 mV (SHE) (El 1 ipsometric film qrowth - jagged trace;

current - solid line trace). From Ambrose and'^Kruger (57).
Fiq. 9, A schematic representation showing the different current

transients expected from straining electrode experiments

and relatina these transients to the amount of material

dissolved after the rupture of a protective film by an emergent slip step. I, rapid reoassivation with little

metal dissolution. II, intermediate repassivation rate

producing restricted dissolution and favoring SCC. Ill,

slow reoassivation producing extensive metal dissolution.

From Staehle flOK

-69—
Fia, 11. A schematic representation of the results obtained in constant

strain-rate tests. Repion A - repassivation is rapid enouah

and strain rate slow enough to prevent damaaina interaction

with the environment. Region B - strain rate high enouah so that reoassi vation does not prevent environriental access. Region C - strain rate is so raoid that ductile nulling apart

takes place.
controlled rate of bare metal exposure, resulting from film rupture by

sl ip 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. 11. For Region A where stain rates are too low to

lead to sec, 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 mechanism (dissolution,

hydrogen embri ttlement, 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 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 (58) 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

see 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 & 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 (23) has made some valuable qualitative connections

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

of the stress intensity factor (described in the next section on mechanics)

but the quantitative relationships require new theoretical and experimental

-72-
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 (61-63).

b) Electron-Microscopy of SCC - With the advent of such powerful tools

as the scanning electron microscope (SEM), and high voltage electron

microscopy, have come new opportunities for the diagnosis of SCC

failures as well as for the prediction of the likelihood of failure.

The predictive ability provided by these tools is based more on the

mechanistic insights they provide (by allowing the examination of the

crack tip itself at nearly atomic resolution) than on any testing

procedure employing them. For example, using mechanisms that rely on

repassi vation as a step in the SCC process, Scully (50) has proposed

that the width of the slip steps produced when a metal is stressed plays

a role in determining SCC susceptibility. He suggests that the wider

the step, the more time is allowed for the migration (under a potential

gradient) of a damaging ion (usually chloride) to the point on the

metal surface where the slip step is emerging. Thus, a greater concentration of this ion will build up at the emergence point and thereby

retard repassivation or enhance dissolution at this point. For fine

sl.ip steps, the time for the build up of concentration of damaging ions

via migration along the emerging step is less and repassivation is more

effective. Therefore, Scully points out that the metallurgical substructure that determines the width of the step becomes crucial . By means of transmission electron microscopy one can determine the metallurgical substructure. For example, it has been found that some alloys

that suffer transgranular SCC have lattices of co-planar arrays of

dislocations (Fig. 12a) which produce wide slip steps [e.g., a brasses

(64)]. On the other hand, alloys where cross slip is easy (Fig. 12b) are

less susceptible to SCC because deformation results from fine slip and

the process described by Scully would be less likely to lead to cracking.

The width of the slip steps produced is probably not critical for all

combinations of environments and alloys and, indeed, has been strongly

questioned (65). Nevertheless, it is mentioned to illustrate how a transmission electron microscopic study of metallurgical substructure

can be used as a failure predicative tool.

Metallurgical substructure plays an important role in the formation

of corrosion tunnels which have been found by some workers using electron

microscopy to be associated with SCC (66,67). Scamans and Swann (68)

describe a technique whereby they are able to look at U-bend specimens

in a high voltage electron microscope and thereby observe the effects

of stress and environmental interactions. High voltage transmission

electron microscopy (500-1000 kV) allows the use of specimens approximately

six times thicker than those used on conventional electron microscopes

(100 kV) and thus come closer to studying phenomena in bulk material.

With double tilting capabilities they could analyze the three dimensional

morpholoay of the crack tip. By this means they could relate attack

morphology of austenitic stainless steels to slip traces and found that more susceptible alloy compositions showed a greater tendency to promote

the formation of corrosion tunnels. They found tunnel diameters increased

-73-
b

Fiq. 12. Transmission electron micrographs showing;

(a) Coplanar array of dislocations leading to wide slipstens

(b) Tangled array of dislocations leadina to fine slipsteps.

(Photographs courtesy of L. K. Ives).

-71;-
with decreasing susceptibility. For the purposes of this paper, a detailed discussion of the role of attaqk morphology and susceptibility

is not necessary. What is important is to stress that by means transmission microscopy one can determine morphological or metallurgical

features (tunnel formation, dislocation array, etc.) which can serve as potential failure predictive aids when their significance has been more firmly established.

Unlike the just mentioned examples of applications of transmission

microscopy, the use of SEM for the study of fractured surfaces

(fractography) is more useful for diagnostic, rather than predictive,

purposes. As such, it is extensively use for failure analysis which

serves as a basis for suggesting measures to prevent future failures.

A comprehensive discussion of the use of fractography in the study of

sec is given by Scully (69). Fig. 13 shows a fracture surface that is

indicative of a failure that was brought about by the hydrogen

embrittlement mechanism of SCC. The flat crysjtal lographic damage

planes are typical of the brittle mode of fracture that would result

from this kind of failure. In contrast, Fig. 14 $hows the dimpled

fracture that results from ductile failure that could not be attributed

to hydrogen embrittlement. The introduction of the SEM in recent years

into many laboratories has given an enormous impetus to fractographic

studies. This is so, because, as pointed out by Scully (69), the SEM

has the following valuable advantages:

1) The microscope has a wide range of magnifications. This is

important in SCC failure analysis because it can quantitatively determine

dimple/cleavage or dimple/intergranular ratios and thereby give

information about the general pattern of crack propagation. Also, by

working at different magnifications it can look at the different varieties

of fracture features that can result from different grain orientations.

2) The microscope has an extremely large depth of field. Because in

SCC where corrosion, corrosion products, and rough surfaces are common,

this feature is of great, actual ly essential , utility.

3) No replication is necessary so that one need not be concerned about

an alteration of the fracture surface (removal of corrosion products,

for example) by the replication process.

4) It can distinguish between the metallic and nonmetallic materials

(corrosion products, inclusions) making up the fracture surface. This can

be a source of misinterpretation when replication techniques are used.

Therefore, the advent of SEM fractography has provided a powerful

tool for the failure diagnosis which is an essential step in the

prevention of SCC failure.
1

Fig. 13. A scanning electron micrograph of the fracture surface

of a high strength steel showing transgranular guasicleavaqe tyoical of hydrogen embrittlement. Arrows show,

dimpled rupture where some ductility is occurring at a few local sites. (Photograph courtesy of C. G. Interrante

and G. E. Hicho).

-75-
Fig. 14. A scanninq electron fractograph showing tyoical ductile

fracture of a high strength steel which indicates no hydrogen embrittlenient. (Photograph courtesy of C. G.

Interrante and R. F.. Hipho),

-77-
Techniques Dealing with the Mechanics of SCC - Up to this point most

of the new techniques described have had their origins in the different

mechanisms of SCC and are aimed at providing predictive abilities.

However, while some offer hope that such opportunities to predict failure

will materialize, the engineer's best present hope lies either in utilizing

past experience or in utilizing laboratory tests that adequately model

service conditions. Presently, the tests developed using the concepts

of linear elastic fracture mechanisms come closest to answering this

latter need. A good review of the theory and application of these

techniques to SCC is given in a book edited by Brown (70). Previously

most laboratory tests for evaluating the susceptibility of an alloy to

SCC in a given environment used smooth specimens subjected to a stress

and measured the time required to observe a failure of the stressed

specimens. One description of the kinds of specimens and experimental

apparatus used can be found in the book by Logan (71). The two major

criticisms of the test methods that do not use the concepts of linear

elastic fracture mechanics are that the geometry of the stress field

around a crack is poorly defined or unknown and that time to failure

involves the time for crack initiation as well as crack propagation.

Brown and Beachem (72) showed, for example, that titanium alloys, which

appeared to be completely immune to SCC because they did not pit readily

and thereby provide initiation sites, were highly susceptible to SCC once initiation sites were present. Therefore, since initiation may be

highly variable depending on whether or not flaws exist in the specimen

studied, the time to failure can be an unreliable and hence unreproducible

measure of susceptibility.

The testing techniques for characterizing SCC susceptibility by the

use of fracture mecham'cs seek to overcome these shortcomings of timeto-failure techniques by providing a pre-existing flaw or crack and by

clearly defining the stress field around a crack through the use of

specimens whose geometry allows the characterization of a parameter K*,

the stress intensity factor. This parameter depends on the loading

and configuration of the specimen, including crack size, and determines

the magnitude of the local stresses. If a certain value of K is found

to cause the growth of a stress corrosion crack, any combination of

structure and crack geometry which duplicates this level of K will

produce the same crack growth rate for the same material and environment.

Thus, using K allows one to relate SCC tests carried out for one crack

and body geometry to another. Brown and Srawley (73) have made

calculations that allow the determination of K for a number of practical

test specimens. The pre-existing flaw used to overcome the problem of

crack initiation is for many specimens a fatigue crack intentionally

introduced for this purpose.

*For a crack of length 2a in an infinite plate where the crack is normal

to the tensile stress field of value a, K = oAa.
The argument for this approach is that most structures have flaws in them

so that crack initiation is not of great practical importance.

Jn sec tests using specimens designed for the application of fracture

mechanics concepts results similar to those shown schematically in Fig.

15 will be obtained. In these tests the rate of crack growth is

measured as a function of K and, as is pictured, there can be regions

(I,II,III) where the functional dependence of crack growth rate on K

differs markedly. Of more importance for SCC testing is the existence

of a threshold value of K, the Kj^^q* shown in Fig. 15, below which the

crack growth rate is zero or insignificant for practical purposes.

Determining ^i^rr important goal in SCC testing because one can relate it to a critical flaw depth, acr^^nd the yield-point stress that

may exist in a structure whose SCC behavior is being determined. Brown

(73) has shown that

= o.2(^^)^ [4] cr ay

where ay is the yield strength. He proposes that a^^ be used as a figure

of merit which includes both Kjsqq and any yield strength stress whether

applied or residual. Fig. 16 shows how one can use a^,^, and ay to evaluate

SCC characteristics of materials. Thus, a material whose value of ^isrr

in a given environment gives a value of X as shown in Fig. 16 would suffer

SCC (crack propagation) if it had a flaw of 0.1 in., but not if it had a flaw of 0.01 in.

There have been a number of criticisms of the general applicability

of the pre-cracked specimens to SCC testing. Parkins (58) has spelled

out a number of these shortcomings. Among these are the following:

1) In order to achieve conditions of plane strain (the conditions

upon which the relationship between K and specimen geometries are based)

for ductile materials, test specimens of impractical size must be used;

2) There may exist problems when pre-cracked specimens are usually

cracked transgranularly although the SCC process for the alloy is

intergranular;

3) The significance of the initiation stage may be to produce a proper

environment for cracking rather than just provide a notch to promote

crack propagation;

4) The frequent observation of crack branching may modify the significance of the values of Kx^-pp determined in fracture mechanics testing.

*The I in the subscript of ^i^cq i^epresents Region I and could be left

out because SCC occurs only in Region I (4).

79-
LOGARITHM OF CRACK GROWTH

RATE

AR£l_T_RARYl iATE

REGION III CRACK GROWTH

RATE STRONGLY DEPENDENT ON

STRESS INTENSITY REGION II CRACK GROWTH

RATE INDEPENDENT OF

STRESS INTENSITY REGION 1 CRACK GROWTH

RATE STRONGLY DEPENDENT ON

STRESS INTENSITY

DESIGNATED

K,.„

STRESS INTENSITY

Fig. 15. Schematic diagram of the data obtained from fracture

mechanics techniques applied to stress corrosion cracking.

From Smith and Piper (74).
YIELD STRENGTH

Fia. 16. A Dlot of Eqn. [4] usina 0.1" a,^6 0.01" value? for a^^

assuminq long surface flaws and yield strength stresifes

onerating. A material with a Kjscc value of X wou'id not

underao SCC in the environment where measurements were

made if surface flaws were no deeper than 0.01". If,

however, the flaws were as deep as 0.1", SCC would occur. From Brown (73)

.
These shortcomings are not serious for the many cases where they

do not apply and where fracture mechanics techniques have been successful,

but it is obvious that one cannot rely exclusively on one testing

approach. Moreover, a recent paper (23) describes ways to interrelate

results obtained from pre-cracked specimens using fracture mechanics

concepts and those obtained from smooth specimens used in constant strain

rate tests. Thus, a combination of the two approaches may be in the

offing.

New Opportunities for Failure Prediction and Diagnosis

This brief survey of new approaches towards developing failure

predictive tools to cope with SCC have identified new opportunities

provided by the three disciplines that govern the phenomenon of SCC.

From chemistry comes the recognition of considering the changes in

environment that occur in the crack and the important effect of environment

on repassivation kinetics. From metallurgy comes the recognition of the

importance of metallurgical substructure and the role that the rate of

bare metal production via slip step emergence plays. From mechanics

comes the recognition of properly characterizing the geometry of the

stress field around a crack. When these tools have been refined

further, both through experiment and theory, the eventual goal of being

able to predict susceptibility to SCC in going from one environment

to another (chemistry), from one metallurgical structure to another

(metallurgy), and from one geometry to another (mechanics) will be

realized. This paper has described some of the beginning steps toward

such a goal

.

ACKNOWLEDGEMENT .

I am most qrateful to the Office of Naval Research which supported

this work under contract NAONR 18-89 NRG 36-082.

On -o2-
REFERENCES

1) R. W. Staehle, Proc. Conf. Fundamental Aspects of Stress Corrosion

Cracking , R. W. Staehle, A. J. Forty, and D. Van Rooyan, eds., Natl.

Ass. Corr. Eng., Houston, 1969, p. 3.

2) B. F. Brown, Naval Research Laboratory Report 7130, June 1970.

3) J. C. Hodge and J. L. Miller, Trans. A.S.M ., 28, 25 (1940).

4) B. F. Brown, "Stress-Corrosion Cracking in High Strength Steels

and in Titanium and Aluminum Alloys," B. F. Brown, ed., Naval

Research Laboratory, U. S. Gov't. Printing Office, Washington,

1972, p. 2.

5) R. W. Staehle, A. J. Forty,and D. Van Rooyan, eds., Proc. Conf . Fundamental Aspects of Stress Corrosion Cracking > Natl. Ass. Corr.

Eng. , Houston, 1969.

6) J. C. Scully, ed.j'The Theory of Stress Corrosion Cracking in Alloys,"

N.A.T.O, Brussels, 1971 .

7) Proc. Intl. Conf. on Stress Corr. Cracking and Hydrogen Embri ttlement

of Iron Base Alloys , St. Etienne, 1973, to be published by Natl.

Ass. Corr. Eng., Houston.

8) B. F; Brown in "Stress Corrosion Crackina of Metals - A State of the

Art,"H. L. Craig, Jr., ed., STP518, A.S^T.M., Philadelphia, 1972.

9) R. N. Parkins, in Ref. (6) p. 167.

10) R. W. Staehle in Ref. (6) p. 223.

11) H. H. Uhlig and E. Cook, Jr., J. Electrochem. Soc , 116 , !73 (1969).

12) E. N. Pugh, W. G. Montague, and A. R. C. Westwood, Corrosion Sci .

,

6, 345 (1966).

13) E. Escalante and J. Kruger, J. Electrochem. Soc . , 118 , 1062 (1971 ).

14) H. H. Uhlig in Ref. (5), p. 86.

15) H. H. Uhlig in Ref. (7).

16) C. M. Preece in Ref. (7).

17) J. R. Ambrose and J. Kruger, Proc. 5th Intl. Cong, on Metallic

Corrosion , Tokyo, 1972 in press.

18) H. L. Logan, J. Res. Natl . Bur, of Stds . , 48, 99 (1952).
19) F. A. Champion, "Intl. Symp. on Stress in Metals and Alloys," Inst,

of Metals, London, 1948, p. 468.

20) J. C. Scully, Corrosion Sci . , 7^, 197 (1967).

21) R. N. Parkins in Ref. (7).

22) T. P. Hoar and J. G. Mines, J. Iron Steel Inst . , 184 , 166 (1956).

23) W. R. Wearmouth, 6. P. Dean, and R. N. Parkins, Corrosion , 29,

251 (1973).

24) D. A. Vermilyea in Ref. (7).

25) A. J. McEvily, Jr. and A. P. Bond, J. Electrochem. Soc . , 112, 638

(1965).

26) E. N. Pugh in Ref. (6) p. 418.

27) K. R. Lawless, Energetics in Metallurgical Phenomena , 1_, M. W. Mueller, ed. , Gordon and Breach, New York, 1965, p. 345.

28) J. Kruger, J. R. Ambrose, and E. Escalante, Natl. Bur. of Stds.

Report 10594, May 1971.

29) N. A. Nielsen, Corrosion , 27. 173 (1971).

30) R. A. Oriani, Ber Bunsen Gesell . f. Phys. Chemie , 76^, 848 (1972).

31) A. R. Troiano, Trans. A.S.M ., 62, 54 (1960).

32) S. S. Birley and D. Tromans, Corrosion , Z7, 63 (1971).

33) M. B. Whiteman and A. R. Troiano, Corrosion , 21, 53 (1965).

34) R. M. Latanision and R. W. Staehle, Scripta Met . , 2, 667 (1968).

35) G. G. Hancock and H. H. Johnson, Trans. Met. Soc. , A.I .M.E . , 236,

513 (1966).

36) J. H. Shively, R. F. Hehemann, and A. R. Troiano, Corrosion , 22^,

253 (1966).

37) M. L. Holzworth and M. R, Louthan, Jr., Corrosion , 24, 110 (1968).

38) D. A. Vaughn, D. I. Phalen, C. L. Peterson, and W. K. Boyd, Corrosion , 19. 315 (1963).

39) Ye. P. Nechay and K. V. Popov, Fiz. Metal Metal loved ., 19, 612 (1965).
B. E. Wilde and C. D. Kim, Corrosion , 28, 350 (1972).

M. Bernstein and H. IL Pickering, Carnegie-Mellon University Report

036-099-2, May 1974.

B. F. Brown, C. T. Fujii, and E. P. Dahlberg, J. Electrochem. Soc .

,

116 , 218 (1969).

G. Sandoz, C. T. Fujii, and B. F. Brown, Corrosion Sci . , 10, 839

(1970).

^

J. A. Smith, M. H. Peterson, and B. F. Brown, Corrosion , 26, 539

(1970).

M. Pourbaix in Ref . (6) p. 17.

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

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

7690, March 1974.

M. A. v. Devanathan and Z. Stachurski , Proc. Roy. Soc . , 270A, 91

(1962).

J. O'M. Bockris in Ref. (7).

J. C. Scully, Corrosion Sci . , 8, 513 (1968).

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

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

T. Murata and R. W. Staehle, Proc. 5th Intl. Cong, on Metallic

Corrosion , Tokyo, 1972, to be published by Natl. Ass. Corr. Eng.,

Houston.

N. D. Tomashov and L. P. Vershinina, Electrochimica Acta , 15 , 501

(1970).

N. Pessal and C. Liu, Electrochimica Acta . 16, 1987 (1971).

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

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

R. N. Parkins in Ref. (6) p. 449.

M. J. Humphries and R. N. Parkins, Corrosion Sci . , 7^, 747 (1967).

D. T. Powell and J. C. Scully, Corrosion , 24, 151 (1968).

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

J., 7, 154 (1972).
62) J. C. Scully and D. T. Powell, Corrosion Sd . . TO,, 371 (1970).

63) D. C. Deegan and B. E. Wilde, Corrosion > 29, 310 (1973).

64) D. Tromans and J. Nuttinn, "Fracture," J. Gilmon, ed., Interscience,

New York, 1963, p. 637.

65) N. A. Nielsen in Ref. (7).

66) N. A. Nielsen, Corrosion , 27_, 5 (1971).

67) P. R. Swann.in Ref. (6) p. 113.

68) G. M. Scamans and P. R. Swann in Ref. (7).

69) J. C. Scully in Ref. (6) p. 127.

70) B. F. Brown, ed., "Stress-Corrosion Cracking in High Strength Steels

and in Titanium and Aluminum Alloys," Naval Research Laboratory,

U. S. Gov't. Printing Office, Washington, 1972.

71) H. L. Logan, "The Stress Corrosion of Metals," John Wiley and Sons,

New York, 1966, pp. 273-285. '

.

72) B. F. Brown and C. D. Beachem, Corrosion Sci . , 5^, 745 (1965).

73) W. F. Brown, Jr. and J. E. Srawley, "Plane Strain Crack Toughness

Testina of Hiah Strength Metallic Materials," A.S.T.M., STP410

(1966).

74) H. R. Smith and D. E. Piper, "Stress-Corrosion Cracking in High

Strength Steels and in Titanium and Aluminum Alloys," B. F. Brown,

ed., Naval Research Laboratory, U. S. Gov't. Printing Office,

Washington, 1972, p. 17.
N2S-1UA

ZZ^T. G- cr.'--* BiBLICCPA=HIC DATA

SHEET

1. iM i'.I.!'. .\ rioN On Kl i'c.'lM' NO.' 2. l..-v't Acct-ssion No. NBSIR 7A-583

3. R«'C ipii iit 's Acn-^ .i.'n

4. 11 i i.i: AMI i< 1 ri l;-. The Role of Passive Filn Growth Kinetics and Properties in

Stress Corrosion and Crevice Corrosion Susceptibility

5. Hublltaiion V'^ic

6. Performing Or^J.ni^ation C-jiic

7. AUTMCR'SI

J. Kruger and J. R. Arirose

8. Perforniinp Or_.in. R.-por: No.

!

9. PERFORMING ORGAMZATiON NAMK AND ADDRESS

NATIONAL BUREAU OF STANDARDS

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3i?n.'. .-

' DEPARTMENT OF COMMCRCE

WASHINGTON, D.C. 20234 NAONR- 18-69. NR 0 36-0'

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

Office of Naval Research

Dept. of the Navy

Arlington, Va. 22217

13. Tvpe of Report i: P<-tioc Cbverei interi":

3/1/73-12/31/73

14. Sponsoring Agency Code

15. SLPPLEMENIARV NOTES

16. A:'. ; iRACT (A 200-i<or7i cr iess tactual -^unmary ol most si&\i(icant intormjtion. If docurrant includes a sigmnc^int

bibliography or literature survey, sjention it here.)

A study of the effect of alloying chroniium and molybdenun in ferritic stainless

steels has been r.ade using techniques to study repassivation kinetics (tribo-ellipsometry) and depassivation kinetics (a recently developed ellips 0:11= try technique for

studying optical changes occurring within a crevice). Results indicate thai: chro-iur.

additions affect both repassivation and depassivation processes x<7hile molybdenun appears

to affect only the repassivation processes. A discussion is given of a generalized

nechanisn for localized corrosion vjhich postulates that susceptibility to attack is

deternined by the competition between depassivation and repassivation processes.

Constant strain rate studies were performed using AISI 304 stainless steel exposed

to acidified l.ON NaCl solution (pH=3.0) . In this particular environment, repassivation

kinetics measurer.ents using the tribo-ellipsometric technique had shown that the rr.etal dissolution rates were rather large due to reduced film repair kinetics, indicating the

possibility of susceptibility to stress corrosion cracking (SCC) . Reduced tir.e to

failure and ma:;inum sustained load at a one strain rate suggest a relationship between

SCC, repassivation kinetics, and rate of pure metal production.

A review of current approaches to the study of stress corrosion and a discussion

of how these approaches can lead to new failure prediction tools is given.

17. KEY 'lORDS (six to t-.velve entries; alphabetical order; capitalize only the iirst letter o( the lirst key word unless a proper name; separated by .semicolons

)

Chloride; crevice corrosion, ellipsometry failure prediction; molybdenum; repassivation

kinetics; stainless steels; stress corrosion; testing techniques

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