NBSIR 78-1429 (Navy)

The Role of Passive Film

Growth Kinetics and Properties

in Stress Corrosion and Crevice

Corrosion Susceptibility

J. Kruger, J. R. Ambrose, J. J. Carroll, and A. J. Melmed

Corrosion and Electrodeposition Section

Metallurgy Division

Institute for Materials Research

National Bureau of Standards

Washington, D C. 20234

November 1977

Technical Summary Report No. 8

Issued May 1 978

Prepared for Office of Naval Research

Department of the Navy

Arlington, Virginia 22217

IMBSIR 78-1429 (Navy)

THE ROLE OF PASSIVE FILM

GROWTH KINETICS AND PROPERTIES

IN STRESS CORROSION AND CREVICE

CORROSION SUSCEPTIBILITY

J. Kruger, J. R. Ambrose, J. J. Carroll, and A. J. Melmed

Corrosion and Electrodeposition Section

Metallurgy Division

Institute for Materials Research

National Bureau of Standards

Washington, D C. 20234

November 1977

DISTRIBUTION OF THIS DOCUMENT IS UNLIMITED

Issued May 1 978

Prepared for Office of Naval Research

Department of the Navy

Arlington, Virginia 22217

U S. DEPARTMENT OF COMMERCE, Juanita M. Kreps, Secretary

Dr. Sidney Harman, Under Secretary

Jordan J. Baruch. Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

PART I

Field Ion Microscopy Studies of the

Interaction of Hydrogen with Selected Metals and Alloys

J. J. Carrol 1 , A. J. Melmed and J. Kruger

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

The entry of hydrogen into a metal's lattice is known to lead to the phenomenon

called hydrogen embrittlement for some alloy systems. It is not known, however,

why some alloys suffer such hydrogen damage and others do not. It is not even

clear which alloys do suffer such damage. For example, it is controversial as

to whether the stress corrosion of austenitic stainless steel is due to hydrogen

entry into the alloy's lattice or to the mechanisms not involving hydrogen.

Moreover, if hydrogen damage is known to occur, its mechanism if, indeed, there

is a single mechanism, is poorly understood and controversial (1).

In order to provide further insights into how hydrogen entry into a lattice produces conditions that lead to sub-critical crack growth, it would

be valuable to examine on an atomic scale if possible, the rearrangements

new phase formations and other disturbances that hydrogen produces when it

enters a metal lattice. The possibility for making such examinations exists

through the use of field ion microscopy (FIM). Recently, this laboratory

completed a study on the effects of hydrogen entry into the lattice of

titanium using FIM (2). The results obtained by FIM for the beginning steps

of hydrogen entry agreed with earlier eletron microscopy studies with regard

to the crystal lographic relationships found for titarium hydride formation.

In the Ti studies the hydrogen was introduced as a gas by using it as an imaging gas. The stress required for hydrogen entry was provided by the high

fields present when using FIM. Hancock and Johnson (3) have shown that in

microscopic systems sub-critical crack growth can be produced in low pressure

hydrogen

The work described here seeks to extend the earlier Ti study by applying

the FIM to an examination of the effect of gaseous hydrogen on a number of

alloys and uranium, a metal known to suffer extensive hydrogen damage. In the group of alloys studied are 1080 steel, 304 stainless steel and Fe-24% Cr

alloy.

EQUIPMENT

The field ion microscope was equipped with a 75 mm diam microchannel plate

image intensifies and a specimen-quick-change chamber for rapid specimen insertion and withdrawal. Routinely, images from two specimens could be obtained

within a half an hour's time. This was an important advantage in the following

studies involving the hydrogen interaction with several types of specimens,

pure metals and alloys, subjected to a variety of environmental and preimaging

treatments promoting hydrogen/metal interactions. Furthermore, the images obtained were generally unstable, thus requiring image intensification in order

to achieve the necessarily fractional second, photographic exposure times.

The specimen temperatures below room temperature were controlled by adjusting

the flow of cold gaseous helium from a liquid helium dewar over specimenmount-feed-through pin. During field ion microscopy reagent grade imaging

1
gases were continuously flowed through the imaging chamber. Imaging gas

mixtures were made by simultaneously opening the appropriate leak valves, and

mixture compositions were estimated from pressure gage measurements.

SPECIMENS

MATERIAL SPECIMEN PREPARATION PROCEDURE

1 . 12 mil diam piano a)

wi re , 1 080 type

,

high strength carbon b)

steel, annealed in

vacuum, 800C, 24 hrs

c)

2. Fe-24%Cr block from a)

which 20 mil diam rods b) were spark cut c)

3. 304 stainless a)

steel , 5 mi 1 diam

wi re

b)

4.

iron whiskers a)

blade shape

b)

5.

depleted uranium, 20 mil a)

diam wire, orthorhombic phase,

annealed 5 days b)

620C

c)

reduce diam and taper, 10% HC10.

i n methanol , 1 5-5Vdc

.

polish to a sharp point, thin

liquid film 1% HC10. in glycerin,

35Vdc while observing at 100X

frequent methanol washes

same as la except 12-2Vdc

same as lb except 2Vdc

frequent methanol washes

reduce diam and polish to a sharp point in 10% HC1CL in

methanol, 50Vdc OR sharpen tip

in 10% HC1 , aqueous thin liquid

film while observing at 100X

water or methanol rinse

use thin liquid film of 1:1:2

HC1 , HNCL, H

? 0, 2-3Vdc OR

10% HCL, aqueous, l-2Vac, and

observe at 100X

frequent water washes

reduce diam and taper using 4 gm

citric acid, 1 ml HNCL, 200 ml

hLO, 40-20Vac J use thin liquid film of same solution to sharpen, 20-25Vac,

while observing at 100X

frequent methanol washes

1080 HIGH STRENGTH CARBON STEEL

Field ion micrographs obtained from this material showed a high

degree of order among the imaging spots indicating cubic symmetry. Fig. 1 shows a field ion micrograph of 1080 steel imaged with a helium-neon gas

mixture. Boundaries (arrows) appeared often in field ion micrographs obtained

from this material. Pending further analysis the boundaries may correspond

either to low angle grain boundaries or twin boundaries, but, for the present,

the boundaries are referred to as grain boundaries. The metal's surface

did not appear to be obscurred by the presence of a surface film. Generally,

disordered alloys and adsorbed metal films tend to image as a random array

of spots. More importantly, interpretation of field ion microscopy studies

of the material's interaction with hydrogen may proceed straightforwardly

without the necessity of considering the influence of an intervening adsorbed

layer

.

2
Exposing 1080 steel to hydrogen gas, with or without the presence

of the electric field, at a specimen temperature of about 30K, did not

have any apparent effect on the steel. Fig. 2 shows a hydrogen field

ion micrograph (no helium or neon present) of the same specimen used in

Fig. 1. More surface atoms were imaged using hydrogen than using heliumneon gas mixtures. The apparent gap between grains was lessened as more

atoms appeared closer to the boundary edge. The crystal planes intersected by the grain boundary retained their definition and integrity,

indicating no adverse interaction in those areas, at least. Fig. 3 shows

another helium-neon field ion micrograph (without hydrogen) of the same specimen immediately after it was subjected to the above described hydrogen

treatment. No adverse effects are apparent. Fig. 3 appears more regular

than Fig. 1 because the imaged area was enlarged by field evaporation

during the hydrogen experiment thereby exposing more step edges to the

imaging electric field.

Exposing titanium and zirconium to similar hydrogen treatments in

the field ion microscope caused micro-cracks and fissures to form in

both metals, and caused crystal planes to appear ruptured (2,4). These

effects were crystallographically specific in each case, and titanium

appeared to be more sensitive to hydrogen than zirconium (4). Clearly,

in the present case of 1080 steel, this did not occur.

A more severe interaction of 1080 steel with hydrogen was observed

at room temperature. Although field ion micrographs show less than optimum

resolution when the specimen is at room temperature, one can see in Fig. 4

a generally gray field corresponding to an adsorbed film, and a darker

central area where the previously adsorbed film was removed by the applied

electric field exposing the 1080 steel substrate. The exposed steel was rapidly field evaporating under these imaging conditions. As a result,

only a couple of net plane rings were photographed but many more were

observed. On the other hand, the adsorbed film appeared stable at best

imaging conditions, with the exception of the interesting bright streaks

seen concentrated in the upper part of Fig. 4. The streaks were unstable

in that they appeared to move about the film surface, and to change in

length. The relatively large ion currents giving rise to the bright lines

may be due to ridge formation on the adsorbed film brought about by the

imaging process. Similarly unstable, bright spotted lines were observed

at cracks formed in titanium observed with hydrogen field ion microscopy

at 25K (2). With correspond!' ngly low temperature imaging applied to the

present case, it may be possible to improve resolution of the room temperature formed film and streaks in order to increase our present understanding

of this interesting feature.

Further film desorption proceeded by exposing the specimen momentarily

to higher applied voltanes than best image voltage. However, in the wake

of the receding film, patches of newly formed film adsorb on the freshly

exposed field evaporating steel substrate. Fig. 5 shows a micrograph of

the same specimen following further film desorption. The crescent shaped

area at the far left shows the remaining adsorbed film of Fig. 4. The

newly adsorbed film patches are brighter than the primary adsorbed film

(crescent area), and are unstable in that the film patches appear and

disappear with time. The dim spotted background image is due to the 1080

3
steel substrate, rapidly field evaporating under imaging conditions.

Fig. 6 shows improved resolution of the same film and substrate as shown

in Fig. 5, brought about by cooling the specimen with liquid nitrogen

cool ant.

The general location of the bright streaks on the room temperature

film shown in Fig. 4 corresponds in part, but not completely, to the

adsorption sites of film patches shown in Fig. 5. It is not known, at

present, whether the film patches of Fig. 5 have the same chemical composition and structure as the adsorbed film of Fig. 4. The damage, if

any, incurred by the 1080 steel substrate as a result of exposure to

hydrogen at room temperature has not been determined.

Additional experiments were attempted in which 1080 steel specimens,

following tip sharpening procedures, were cathodically polarized between

5 and 10 volts in an aqueous 10% HC1 0 solution for a few seconds in order

to increase the amount of hydrogen incorporated into the metal lattice.

Cathodic charging adversely affected a specimen's lifetime, that is all

specimens prematurely failed during field ion microscopy. No specimen

survived neon field ion evaporation and imaging conditions long enough

to obtain a crystallographic map of the steel substrate. However, room temperature hydrogen imaging produced similar films as those in Figs. 4

- 6. Further studies are necessary to establish the effect that hydrogen

may have on 1080 steel lattice and grain boundaries that are visible with

field ion microscopy.

304 STAINLESS STEEL

Generally, 304 stainless steel did not image as well as 1080 steel

in that most spots in the field ion micrographs appeared to be randomly

distributed. From the few net plane rings that did appear about the

surface, there was not enough information to crystal 1 ographical ly index

micrographs. For example, some net plane rings appeared in the central

grain of Fig. 7 (arrows). The central grain was roughly elliptical in cross section and distinguished from the other grain by a dark, finite

width line. During the course of field evaporation the central grain

became smaller in size and finally disappeared. In summary, there is

a modicum of ordered spots in a typical field ion micrograph of stainless steel, and evidence that grain boundaries are visible.

Stainless steel /hydrogen interaction experiments were carried out

at 30K. No changes were detected in the appearance of the pertinent

field ion micrographs. The reason for this may be due to the presence

of an adsorbed, contamination layer (giving rise to micrographs showing

little overall crystallographic order) interfering or altering the

hydrogen/stainl ess steel interaction that would occur otherwise.

IRON - 24% CHROMIUM

Fig. 8 shows a neon field ion micrograph obtained from an Fe24% Cr specimen. Field ion micrographs obtained from Fe - 24% Cr and

304 stainless steel (Fias. 7 and 8) show similar imaging characteristics , - only a few net crystal plane rings visible in a field of disordered spots.

Exposing and imaging with hydrogen gas produced completely disordered

field ion micrographs (Fig. 9) indicating that these imaging conditions

may have produced an adsorbed film.

4
Evidence was obtained that suggests that the images of Figs. 8 and

9 came from an adsorbed film, at least in part. Fig. 10 is a field ion

micrograph obtained from the same specimen that was used in Figs. 8 and 9,

but imaged in a peculiar way that deviates from usual field ion microscopy practice. The specimen was subjected to a several thousand volt

pulse over the voltage required for imaging. During the pulse the specimen

surface was dramatically destabilized and several hundred layers were

removed quickly by the field evaporation process. Stabilizing the image

again at best image voltage. Fig. 10 was obtained. The pattern is regular

enough to be crystal lographical ly indexed. The specimen may be clean

enough to permit examination of its hydrogen interaction in more detail.

Furthermore, the pulse field evaporation technique applied to 304 stainless

steel may produce improved field ion micrographs.

IRON

It is known (5) that pure iron field ion micrographs show excellent

overall order and atomic detail on the (110), (100), (112) and higher

order crystal planes, and with special care taken to exclude minute quantities of hydrogen in the gas phase, the (111) crystal plane as well. With

normal background partial pressures of hydrogen in unbaked, ion pumped,

metal vacuum systems, the (111) plane area images do not show net plane

rings among the imaging spots (5). See Fig. 11. The (111) plane area

of iron appears to be more sensitive to hydrogen than other crystal planes.

Flydrogen field ion micrographs of iron obtained at 30K, as shown in Fig. 12,

retained net (110) and (100) crystal plane order but the (111) plane area became move disordered and imaged more brightly. However, hydrogen imaging

conditions at 30K did not have any lasting effect on iron, for when neon

was substituted for hydrogen as the imaging gas, the (111) plane areas

imaged first and quickly ordered during field evaporation as the rest of

the crystal areas imaged in ordered fashion. No evidence in the iron field

ion micrographs was found that indicated defective or irregularly shaped

crystal planes, micro-crack defects of adsorbed films.

Severe corrosion of iron occurred, however, when iron was exposed

to trace amounts of air during neon field ion microscopy. Fig. 13 shows

an iron field ion micrograph obtained under these conditions. No spots

were observed in the (111) plane area, indicating that a hole or pit

was formed in this area, perhaps, as a result of preferential field evaporation of reaction products at significantly less field strengths than

that required for neon imaging.

URANIUM

Figure 14 shows a field ion micrograph obtained from uranium. Until

now, no orthorhombical ly structured element has been successfully imaged

in the field ion microscope. The image in Fig. 14 was obtained from a single

crystalline grain. The image contains many crystal planes continuously distributed about adjoining and symmetrically disposed bright and dark contrast

field regions. The problem of crystal lographical ly indexing uranium field

ion micrographs has not been resolved. Presently, the grain size of the source material is too small to obtain single crystalline x-ray diffraction

patterns

.

5
After prolonged imaging using hydrogen, patterns typified by Fig. 15 were obtained. The pattern shows several grain boundaries and irregularly

shaped grains ranging between about 5-30 nm in breadth. After replacing

hydrogen imaging gas with neon, followed by considerable field evaporation,

it was found that the small grained surface image was followed by a uniformly disordered image that, upon further field evaporation, showed film

desorption characteristics. As the disordered film desorbed an image with

similar characteristics as Fig. 14 was exposed. The evidence indicates

that the micrograph in Fig. 15 was obtained from a surface hydride of

uranium. Supporting evidence was found by exposing a uranium specimen,

showing neon images similar to Fig. 14, to hydrogen field ion imaging

conditions. A disordered film formed immediately, followed by the gradual

formation of randomly scattered small clusters of bright spots, the number

and size of the clusters increasing with hydrogen exposure time. Figs.

16 (a , b ) were obtained during this experiment after about 15 and 30 minutes

of hydrogen field ion microscopy, respectively. These field ion microscopy

experiments indicate that, in the case of uranium in the presence of

hydrogen field ion imaging conditions, an adsorbed layer was formed followed

by nucleation and growth of uranium hydride.

CONCLUSIONS

In the foregoing studies an attempt was made to broaden our current

knowledge of the interaction of hydrogen with pure metals and alloy steels

by using the technique of field ion microscopy. Through the application

of this technique many microscopic details were observed under the ex- perimental conditions imposed by the technique, namely hydrogen adsorption on field evaporated surfaces in the presence of large electric fields.

The following list includes the presently obtained results. It is apparent

however, that further work is necessary in order to complete the investigations .

1. Mo evidence was found that indicated that 1080 steel was effected

adversely by its interaction with hydrogen between 30-80K.

2. Although further experimentaion is necessary before the situation

at room temperature is clarified regarding possible 1080 steel substrate

damage brought about by its interaction with hydrogen, it was evident

that substantial film adsorption occurred on specimens with or without

a previous history of cathodic charging.

3. Control experiments with pure iron indicated that a) the (111) plane

area was most sensitive to hydrogen, as evidenced by a damaged surface

layer, and 2) the (110) and (100) planes appeared uneffected.

4. Experiments with 304 stainless steel and iron-24% chromium steel

indicate that ordinary imaging techniques may not reveal film-free,

clean metal surfaces. No significant effects were observed using such

surfaces in hydrogen experiments. A new technique was devised that indicated

6
that iron-24% chromium can be cleaned from the strongly adsorbing film

ordinarily observed. Further work using the revised technique is necessary

in order to determine the technique's usefulness in obtaining cleaner

stainless steel surfaces. Hydrogen experiments involving cleaner surfaces of

iron-24% chromium and 304 stainless steel may show interesting effects that

may reveal the type of interaction involved, kinetic information, and the

role of alloying additives.

5. Hydrogen experiments with uranium indicate that hydrides readily form

by a process of nucleation and growth on a previously formed adsorbed layer.
REFERENCES

1. R.W. Staehle, "The Theory of Stress Corrosion Cracking in Alloys"

NATO Scientific Affairs Dir., Brussels, 1971 pp 231 - 243.

2. J.J. Carroll, J. Smit and A.J. Melmed, "Fiel ding-ion imaging of titanium

and effects due to hydrogen," Surface Sci. 45 (1974) 640 - 648.

3. G.G. Hancock and H.H. Johnson, "Hydrogen, oxygen, and subcritical

crack growth in a high-strength steel," Trans Met. Soc. AIME 236

(1966) 513 - 516.

4. J.J. Carroll and A.J. Melmed, "Field ion microscopy of zirconium,"

Surface Sci. 58 (1976) 601 - 604.

5. Erwin W. Muller and Tien Tzou Tsong, Field Ion Microscopy; Principles

and Applications (Elsevier, New York, 1969), pp. 116, 161 and 174.

8
Fig. 1. A field ion micrograph obtained from a 1080 steel specimen, imaged

in a gas mixture of 55% helium and 45% neon, at a specimen temperature of

25K, 3750 volts.

Fig. 2. A hydrogen field ion micrograph obtained from the same 1080 steel

specimen as imaged in Fig. 1, 25K, 2600 volts. Note the appearance of two

dark areas near the center of the micrograph. Both are artifacts that appear in most of the following micrographs.

9
Fig. 3. A field ion micrograph obtained from the same 1080 steel specimen

as imaged in Fig. 1, following its exposure to the hydrogen imaging con- ditions described in Fig. 2, and imaged in a 55% helium and 45% neon gas

mixture, 25K, 5000 volts.

Fig. 4. A hydrogen field ion micrograph obtained from a at room temperature, 8800 volts. See text for additional

080 steel

details.

specimen

10
Fig. 5. A hydrogen field ion micrograph obtained from a 1080 steel specimen

at room temperature, 8200 volts.

Fig. 6. A hydrogen field ion micrograph obtained from the same specimen

as imaged in Fig. 5, but cooled with liquid nitrogen, 11,000 volts.

11
Fig. 7. A neon field ion micrograph obtained from a 304 stainless steel

specimen, 30K, 12,500 volts. Arrows indicate crystal net plane rings.

Fig. 8. A neon field ion micrograph obtained from a specimen of iron - 24% chromium, cooled at 30K, 14,000 volts.

12
Fig. 9. A hydrogen field ion micrograph obtained from the same specimen

as imaged in Fig. 8, 30K, 10,000 volts.

Fig. 10. A field ion micrograph obtained from the same specimen as imaged

in Fig. 8, imaged in a neon and hydrogen gas mixture, 30K, 15,000 volts.

i I

13
Fig. 11. A neon field ion micrograph obtained from iron, 30K, 8600 volts.

Fig. 12. A hydrogen field

specimen as used in Fig. 11 ion micrograph obtained from the same iron

, 30K, 6600 volts.

14
Fig. 13. A neon field ion micrograph obtained from an iron specimen

exposed to a small partial pressure of air, 30K, 9400 volts.

15
Fig. 15. A hydrogen field ion micrograph obtained from uranium, 30K,

8600 volts.

16(a)

30K

14,700 volts

16(b)

30K

16,200 volts

Fig. 16 (a, b) Flydrogen field ion micrographs obtained from a uranium specimen exposed to hydrogen imaging conditions for about 15 and 30 minutes,

respecti vely.

16
PART II To be published in Proceedings of 4th

International Symposium on Passivity

The Role of Noncrystall ine Films in

Passivation and Breakdown of Passivation

A. G. Revesz

COMSAT Laboratories

Clarksburg, MD 20734

and

J. Kruger

National Bureau of Standards

Institute for Materials Research

Washington, D.C. 20234

1 . INTRODUCTION

In his 1969 Palladium Medal address. Hoar 1 set forth some "rules"

to serve as guidelines for producing alloys that would exhibit the best

corrosion resistance because of their ability to form superior passive

films. One guideline of his is to use metals in an alloy that lead

to the formation of "mixed" oxides involving polyanions which result

in "cross-linked monolithic amorphous structures not containing the

extremely defective grain boundary regions of polycrystall i ne materials". Due to the lack of grain boundaries in such a noncrystall ine

film, the outward migration of cations is reduced.

Subsequent to this suggestion by Hoar, there were a number of

experiments and discussions that have served to bolster his insight.

For example, McBee and Kruger 2 showed by electron diffraction that

as one adds Cr to Fe in an alloy the films formed go from polycrystalline to noncrystal 1 i ne structure with increasing Cr content.

According to Okamoto 3 "the most important parameter controlling the

corrosion resistance of [stainless! steel is concluded to be the

amorphous nature of the film in which bound water is included."

Leach 4 has also noted that glassy films tend to promote corrosion

protection in many cases. He also pointed out that this enhanced

corrosion resistance correlates with the greater ductility of the

oxide film. Fehlner and Mott, 5 while not being concerned about

passivity, asserted in a review on low-temperature oxidation that

"an oxide film formed at low temperature is amorphous." They

have suggested that, depending on whether the oxide is a "network

former" or "network modifier" (using Zachar-iasen 1 s 6 criteria)*,

the amorphous film is with or without a network structure, and

that this structural difference determines the transport process.

However, details of the noncrystal 1 i ne structure, particularly in

the case of network modifiers, were not discussed. Lawless 7 has

pointed out in his review that the low temperature oxide films on

Cu, Ni , Zn, and Fe are known to be crystalline. Some aspects of

Fehlner and Mott's assertion concerning the formation of

amorphous oxide film on aluminum have been critized by Pryor. 8

*An important feature of the glass network is that oxygen atoms

(ions) form M-O-M bridges, i.e., their coordination is two.

I
Judging from the examples just described and others, it appears

that there is increasing evidence that an important aspect of the

beneficial properties of some passive films are in many instances

related to their noncrystall i ne structure. From the viewpoint of

passivation, very little attention has thus far been given to a detailed examination of what are the structural characteri sti cs of noncrystalline oxide films on metals. Are there different

types of noncrystall ini ty? What factors influence the structure

and stability of noncrystall i ne films; what effect do these

factors have on passivity and on breakdown of passivity?

These questions are also very pertinent to some other areas outside passivation. Closely related to passivation is the

formation and properties of anodic oxide films, many of which are noncrystalline (e.g. Ta 2 0 5

, A1 2 0 5 ).

9 The importance of non- crystallinity for oxides used as dielectric films in capacitors

has been pointed out by Harrop and Campbell 10 but without

discussing the fundamental problems of structure.

Some fundamental questions have been discussed from 1967 on in

connection with Si0 2 films on silicon. 11 This film is used to

stabilize the surface properties of silicon devices; the

discovery in 1959 by Atalla et al„ 12 that thermal oxidation of

silicon stabilizes, i.e. passivates, its surface, was a crucial

step in semiconductor device technology. Revesz 13 has reviewed

the surface passivation of silicon emphasizing the importance of

noncrystallinity.* More recently, the work has been extended to

noncrystall ine Ta 2 0 5 films on silicon. 14

, 15 In the course of

these studies, concepts and results developed that appear to be

pertinent to understanding the role of noncrystal 1 ini ty in

influencing the properties of passive films on metals.

From the classic point of view of passivity, films on silicon may

not seem relevant. However, since the S i / S i

0

2 interface is the

best understood solid/solid interface and S i

0

2 is a non- crystalline (vitreous) solid par excellence , the passivation and

oxidation behavior of silicon can be used as a basis of

comparison in discussing more conventional systems such as stainless steel. It is, thus, the plan of this paper to develop

a comprehensive discussion of noncrystall ini ty using the oxide

films of Si and Ta as a basis and then to discuss the role that

noncrystallinity can play in passivity and the breakdown in

passivity in general.

*The development of integrated circuits and, thus, today's computer

industry is largely due to the excellent properties of vitreous Si

0

2 film

and its interface with the Si substrate. This has been a very sophisticated

applications of semiconductor surface passivation which is very close indeed

to passivation of metals.

18
2. NONCRYSTALLINE OXIDE FILMS AMD FILM/SUBSTRATE INTERFACES

In this section, various aspects of noncrystalline oxide films

and film/substrate interfaces will be discussed using Si02 and Ta 2 0 5 films as examples. These two oxide films are used to demonstrate the

importance of bond and structural flexibility in noncrystal 1 i ne structures. The usefulness of these concepts will be illustrated on a few other oxides which are important from the viewpoint of metal

passivation.

2 . 1 The Si/SiOo Interface Structure

Silicon single crystal is the most perfect and best understood

solid. Owing to its great importance in semiconductor device

technology, the Si/Si02 interface has been extensively studied. An

important factor in these investigations has been that the

noncrystall i ne Si0 2 films are prepared on a very perfect substrate

under very clean conditions. As a result, the Si02 films are also

quite perfect. The electronic properties of the Si substrate are very

sensitive indicators of the perfection of the Si0 2 film, including the

Si/Si02 interface, unsurpassed by any other physicalchemical

analytical technique. This is why the Si/Si02 interface has emerged

as the best characterized solid/solid interface. This section is

essentially based on a review, 13 and works listed in the bibliography

there are not referenced here.

Silicon is a very electropositive element but, in practice, it

is inert in aqueous media due to the passivating oxide film. This

noncrystal 1 i ne oxide film forms at a logarithmic rate in room temperature conditions reaching a thickness of about 3-^. 16 This

behavior is similar, among others, to that of aluminum or tantalum;

but, in contrast to these metals, thermal oxidation of silicon (even

at very high temperatures) results in a noncrystall ine Si02 film. The

perfection of thermally grown Si02 films can be illustrated by the low

density of sodium impurity (<10 15 cm

-3

) and electron trapping defects

(<10 14 cm- 3 ), high dielectric breakdown strength (>106 V cnr 1 ),

excellent reproducibility, and stability under extreme conditions.

The perfection of the Si/Si02 interface is demonstrated by

the very low density of interface states: <^10 10 cm- 2 (corresponding

to ^ 10- 5 monolayer adsorbate). These interface states are associated with unsaturated ("dangling") bonds of surface Si atoms

(which at an ideal Si/Si02 interface should be bonded to oxygen

atoms) and/or with electron (hole) trapping defects in the oxide

film within tunneling distance (<^15A) from the interface. These

interface states may act as localized traps for electronic (holes)

in the silicon and, thus, adversely affect its electronic properties.

Also, they represent chemically active sites and, as such, they were responsible for the surface instability effects which plagued semiconductor devices before the discovery of surface passivation. In

19
this latter sense, interface states at the S i / S i

0

2 interface are,

to some extent, similar to active sites in metal/passivating film

interface structures which play such an important role in the

breakdown of passivation.

The high perfection of the Si0 2 film and of the S i / S i

0

2 interface is due to the noncrystal 1 i ne structure of S i

0

2

- When the film

begins to crystallize (e.g., during prolonged heat treatment at high

temperature) the properties of the film (e.g. dielectric strength)

deteriorate and the density of interface states increases. The

first effect results from the disruption of the Si-0 network as crystallites form. This process is associated with the generation

of network defects and, later, even with the formation of grain

boundaries between crystallites. As was already mentioned in the

Introduction, the presence of grain boundaries in the oxide film

leads to the loss of its passivating (protective) character. Thus,

a polycrystal 1 i ne or highly defective noncrystalline oxide film is

not a good candidate for passivation.

The increase in interface state density during crystallization

of the S i

0

2 film results from the loss of structural flexibility

characteristic of the Si0 2 film while it is noncrystalline. As a result of this flexibility, the majority of the surface Si atoms

can be bonded to an oxygen atom of the oxide without requiring an almost perfect epitaxial relationship between Si and S i

0

2

* When,

however, the Si

0

2 film is crystalline, the match of two structures

with rigidly defined but different long-range order is poor and,

hence, the interface is disordered. Apparently, quasi -perfect

solid/solid interfaces can be only obtained when either a quasiperfect epitaxial relationship exists between two single crystals

or at least one side of the interface is noncrystalline. This

conclusion is obviously very pertinent to the passivation of

metals which are usually (but not exclusively) polycrystalline and,

therefore, an epitaxial relationship with the oxide film is out of

the question. Thus, elimination or drastic reduction of the active

sites in the thin passive film requires that the passivating film

be noncrystal line .

It can be seen from the above discussion that Si

0

2 films can

be very perfect in spite of, rather, because of the fact that they

are noncrystalline. The reason for this is that the lack of longrange order does not simply arise from a destruction of structural

order altogether. Rather, the short-range-order characteri Stic

of crystalline polymorphs of Si0 2 is maintained in noncrystalline

Si0 2 films, as well as in fused silica; that is, every Si atom is

surrounded tetrahedrally by four 0 atoms and the Si-0 bond length

is constant. It is the Si-O-Si bridging bond angle which is

characteri zed by a wide distribution (from ^120° to 180°) in

vitreous Si0 2. Because of this distribution it is possible to

establish a Si-0 network without long range order and without

breaking many Si-0 bonds and, hence, introducing a large number

of defects. This type of noncrystalline structure has been named

20
vi treous in contrast to amorphous noncrystall ine structures which

do not have appreciable short-range order. 11 Fused silica and Si

0

2 films on silicon are vitreous solids par excellence , whereas, for

instance, vacuum deposited SiO is an amorphous solid. The lack

of short-range order in amorphous solids is usually associated

with chemical disorder, whereas the short-range order in vitreous

solids ensures perfect stoichiometry and excellent reproduci bi 1 i ty

.

The structural flexibility of Si

0

2 is also manifested in the

large number (9) of crystalline polymorphs which have essentially

identical short-range order and almost the same standard free

enthalpy of formation (the difference between quartz and fused Si0 2 is about 1 percent); yet, they exhibit significant variations in

several properties, e.g. , density, bond polarizabi 1 ity , etc.* The

reason for this behavior is that various topological arrangements

of SiO u / 9 tetrahedra of essentially the same energy can be present

in these polymorphs. Thus, similar to polymers, the variation in structure

is essentially a conformational problem .

2.2 Bond Flexibility and Noncrystalline (Vitreous) Structure

The structural flexibility of Si

0

2 polymorphs, including

vitreous Si0 2 is due to the flexibility of the Si-0 bond. This bond

is partially ionic (%40%) and partially covalent (^60%) and the

covalent part has a and it components.** Apparently, the individual

contributions of these components can vary without significantly

affecting the total bond energy but influencing other properties.

On this basis, we can visualize vitreous Si

0

2 as a network built

up from S i

0

4 / 2 units where a large number of local conformations

of almost the same energy exist and, accordingly, the Si -0-Si

angle varies greatly. The flexibility of the Si-0 bond is thus

the principal reason that SiO2 can be prepared in the vitreous

state.

The importance of partially ionic and partially covalent

bonds in vitreous structures has been discussed in great detail

by Myuller 17

. However, the question of bond flexibility is more complicated than the considerations of the mixed ionic-covalent

bond indicate. This point is illustrated by a comparison between

Si

0

2 and Ge0 2

.

*With the exception of al imrinLun, no metal has a comparable number

of oxide polymorphs. Aluminum oxide (corundum) has five polymorphs whose

standard free enthalpies of formation are within 2.5% of each other.

**The o-bond arises from sharing one of the 3sp 3 electrons of silicon

with on of the 2p valence electrons of oxygen; the ir-bond is due to the

overlap between the 3d orbitals of Si (empty in the ground state) and the

2p orbital of 0 containing the lone pair electrons.

21
In spite of the close resemblance of germanium to silicon,

the properties of Ge0 2 are rather different from those of Si0 2 , even though the Ge-0 bond is also a mixed ionic covalent bond,

roughly in the same proportion as the Si-0 bond. However, there

is no iT-bonding in Ge02 because of much less overlap between

the oxygen 2p and germanium 4d orbitals relative to the Si 3d

orbitals. As a result, the bond flexibility is greatly reduced.

Thus, there is only one Ge02 crystalline polymorph with 4:2

coordination (hexagonal, silica structure) and the tendency

toward noncrystal 1 i ni ty is much less than in the case of Si02

. The

lack of TT-bonding in Ge02 is also partially responsible for the

significant solubility of hexagonal and noncrystalline Ge0 2 in

water. Hence, the thin Ge0 2 film formed at room temperature does

not passivate to such an extent as the Si02 film does, and

thermally grown Ge0 2 film is useless for surface passivation

since it is soluble in water and chemically very reactive.

This difference in the passivation behavior is the main reason

that silicon has replaced germanium in semiconductor device

technology.

A different kind of bond flexibility is exhibited by metals

which can be in various oxidation states. Thus, for example the

standard free enthalpies per g-atom of Cr are -141 and -136 kcal mole- 1 for Cr0 3 and Cr2 0 3

, respectively. This difference is relatively small

so that various structural conformations of roughly the same energy

may exist. This could be a factor in the formation of relatively

stable noncrystalline In addition, transition metal oxides, e.g. Cr2 0 3 , may exhibit a departure from the regular structure due to ligand

effects. 18 Since both o- and irbonding can be involved in these

effects and ir-bonding can affect the stereochemistry while its

contribution to the bond energy is relatively small, a situation

somewhat similar to that characteristic of the Si-0 bond may arise.*

Even without involving ir-bonding, ligand effects may increase the bond

flexibility. It is possible that these effects are partially

responsible for the important role of H 2 0 in passivation since H 2 0 is

a a donor ligand. The effects of H 2 0 will be discussed below, but it

is significant to mention in this context that due to the large number

of delectrons, chromium (electron configuration in ground state : 3d 5 4s 1

) can be associated with water as a ligand in several

configurations; this might be an additional important factor in

increasing the stability of the noncrystalline passivating film on Cr

metal and in the beneficial effects of Cr alloying on corrosion.

2.3 Vitreous Ta^Oc Films and Structural Flexibility

It is well known that anodic oxidation of tantalum results

in a noncrystalline Ta 2 0 5 film whose dielectric and optical properties

as well as its excellent reproducibility, indicate that there is a

*The 0

2 -ion as a ligand can act as o or n donor. In the Si-0

bond the role of oxygen can be considered as similar to that of rr donor ligand.

22
high degree of short-range order in the structure, i.e., the film

is vitreous. 9 On the other hand, thermal oxidation of tantalum,

both in polycrystalline and single crystalline form, usually

results in polycrystalline oxides with poorly defined properties. 8

, 19

Apparently, the thin noncrystall i ne passivating film present on

the Ta surface breaks down during thermal oxidation, whereas it

grows further during anodic oxidation. Revesz and coworkers 14 have found that vitreous Ta 2 0 5 film can also be obtained by

thermal oxidation of Ta film vacuum deposited on silicon. This

Ta 2 0 5 film is morphologically and optically so perfect that it is

used as an anti refl ection film in silicon solar cells of improved

conversion efficiency. This vitreous Ta 2 0 5 film on silicon can withstand prolonged heat treatments without crystallization.

Polycrystalline Ta foil oxidized under identical conditions

resulted in polycrystalline Ta 2 0 5 film of very poor optical

properties. These observations demonstrate that the substrate

plays a great role in determining the formation and stability

of a noncrystal 1 ine (vitreous) solid. This aspect will be

discussed below, but the fundamental question is: why can Ta 2 0 5 be obtained under some circumstances in the vitreous state?

In order to answer this question, we have to consider the

structure of crystalline Ta 2 0 5

. According to Stephenson and

Roth, 20 the structure of Ta 2 0 5 is based on chains of distorted

octahedral and pentagonal bipyramidal oxygen polyhedra around the

Ta atoms. The coordination of oxygen in these configurations is

not two; hence, simple Ta-O-Ta bridges (analog to Si-O-Si

bridges) do not exist. Consequently, the Zachariasen rules of

glass formation 6 are not fulfilled and, in this sense Ta 2 0 5 cannot be considered as a "network former". However, it is

network former on the basis of metal -oxygen bond strength. 5

Since the Ta-0 bond is ^60 percent ionic and ^40 percent

covalent, the bond is expected to have some flexibility, even

though, as was shown for the Ge-0 bond, this kind of flexibility

is not necessarily sufficient. Additional bond flexibility may

arise from ligand field effects characteristic of transitionmetal oxides as mentioned above. This is probably the reason for

the large variation (from 1.83A to 2.60A) of the Ta-0 bond length

in crystalline L-Ta 2 0 5 (the low temperature polymorph). 20 The

structural complexity of this structure is manifested in the very

large unit cell that is divided in subcells. 21 The arrangement of

these subcells and the size of the unit cell depend on heat treatment and impurities such as hydrogen, silicon, etc.; for example,

the unit cell formula at temperatures below-1000°C is Ta 28 0 70 .

The structural complexity is closely associated with the possibility of

varying the composition (stoichiometry) by changing the relative numbers

of structural subunits which are structurally closely related but differ

slightly in their composition; these structures have been named "infinitely

adaptive" 22 or "vernier" 23 structures. Since the difference in lattice

23
energy between the alternative structures arrangements is very low, the

various substructures can easily interchange. In this regard, the structure

of L-Ta 2 0 5 exhibits a similarity to polymers and, analogously, the various

arrangements of substructures might be considered as structural confor - mati ons .

It is very likely that the adaptive nature of the L-Ta 2 0 5 structure

plays an important role in the formation of noncrystalline Ta 2 0 5. Specifically, we suggest that the increased tendency toward longrange order usually exhibited by structures in which the nonmetallic

element has a coordination larger than two is counteracted to some extent by the structural complexity in Ta 2 0 5

. As a result, arranging

the atoms according to the long-range order of the crystal during the

formation of the oxide may be more difficult kinetically than forming

a vitreous structure.

Since an interaction with the Si substrate occurs during the

oxidation of Ta film on silicon , 15 it is interesting to note that

incorporation of silicon leads to increased complexity in the

unit cell of L-Ta 2 0 5

. 21 As a result of this increased

complexity, silicon hinders the formation of the crystalline

structure. Consequently, silicon facilitates the formation of a noncrystall i ne structure and may also stabilize it.

Since the coordination of A1 and 0 atoms in A1 2 0 3 is 6:3, the

oxygen does not form A1-0-A1 bridges and so this oxide also cannot

strictly be considered as a network former. However, it was mentioned

above that there are five polymorphs of AI 2 O 3 of not too different

formation energies. The structural difference among these polymorphs

arises from different degrees of ordering A1 atoms in an essentially

closest packing of 0 ions . 24 This structural flexibility is quite

different from that exhibited by Si0 2 but it could be important in

the formation of noncrystalline A1 2 0 3 as well as in other ionic oxide

structures based on nearly closest packing of 0 ions. Such an oxide is,

for instance, the Fe-Cr perovskite, which is the structure of oxide

films on stainless steel with low Cr-content . 2 It is very likely

that the structural flexibility of the perovskite structure increased

by the incorporation of chromium and this could be the reason that

at higher Cr content the passivating film is noncrystall ine.

Therefore it is evident that the most important factor deter - mining whether an oxide film can be obtained in a reasonably stable

vitreous form (i.e., in a noncrystall ine structure with a high degree

of short-range order) so that it can passivate the underlying metal

is that it must have some flexibility . This flexible structure then

allows the formation of localized structural conformations having

nearly the same energy. These structural units are essentially

identical to those present in the crystal but they are linked

together without long range order. This flexibility is more

decisive than the geometric requirements. Clearly, classification

of oxides into network formers and network modifiers is not

sufficient. The problem of noncrystalline (vitreous) structure is

24
fundamentally associated with the bond/structure flexibility but

many factors influencing the kinetics of formation and stability

are involved . Essentially, the problem lies in kinetics rather than

theromodynamics.

3. GENERAL CONSIDERATIONS OF FACTORS INFLUENCING NONCRYSTALLINE

STRUCTURE

It is evident from the preceding discussion that a passivating

film should be noncrystalline, preferably, vitreous in the sense of

having appreciable short-range order. The intrinsic structural flexibility is obviously an important requirement but this can be modified

for instance, by incorporating other atoms in the oxide film, as well

as by various factors which influence the kinetics of the formation,

self-repair of localized break-down, and formation of passivating

noncrystal 1 ine films. Some of these factors have already been

briefly mentioned in connection with the oxide films discussed above

and are systematically discussed in this section.

3.1 Effects of Additives to the Oxide Film

Concerning the corrosion resistance of alloys. Hoar 1 has con- cluded that "we may expect mixed or compound oxides on alloys usually

to be more protective than on pure metals because they will tend to be

less soluble and less 1 crystal 1 ine' " . Additives to the oxide film can

be introduced not only by alloying the metal but by surface treatments

e.g., ion implantation, as reviewed recently by Ashworth et al. 25 In view of the discussion in the preceding section, it is obviously

desirable that the additive enhances noncrystall inity of the passivating

film. This can be achieved by increasing the structural/bond flexibility

as for instance, Cr does in the passivating film on stainless steel.

Another possibility is to increase the stability of the noncrystalline

structure by decreasing the ease of the crystallization; i ncorporation

of Si into Ta 2 0 5 , mentioned above, is an example.

Special attention should be given to hydrogen as an "additive" in

oxide films. Hydrogen in the oxide film can lead to increased

structural flexibility by forming M-OH bonds in addition to M-0

bonds. If the correspond!’ ng metal hydroxide has various

polymorphs as, for instance, A10.0H, 26 then various conformations

of the M-OH bond may exist in the H- containing passivating film;

this results in increased tendency toward noncrystallinity. The

M-OH groups can be linked together via hydrogen bonds as in the case of A10-0H and CrO.OH; 26 this facilitates the formation of a noncrystalline network (cf . H-bonding in organic polymers).

However, H or OH may also weaken the oxide structure in some

cases. The fact, that hydroxide ions of some metals (e.g. Cr,

Al , Ta, Si) can form hydrogen-bonded polymer (but not Fe, Ni , Cu,

etc.) 27 is probably also pertinent to the structural flexibility

of noncrystall ine films.

Another important property of hydrogen (water) in passivating

films is its role in self-repair as has been pointed out by Okamoto. 3

25
This effect is somewhat analogous to the reduction of interface states

(dangling Si bonds) at the Si/Si02 interface by forming Si -H or Si -Of! bonds. 13 In this manner the reactive sites can be tied up so that the

apparent perfection of the Si/Si0 2 interface is increased. However,

this interface is more sensitive to destructive effects (e.g. irradiation) than a quasi -perfect S i / S i

0

2 interface where practically all

Si atoms are bonded to oxygen.

Hydrogen may be present even in films grown at high temperatures

under relatively clean conditions: for example, Si0 2 films or silicon

may contain S i OH and/or SiH groups up to 1022 cm- 3

. 28 Evidently, a very

high concentration of hydrogen is to be expected in thin passivating

films formed at room or in aqueous environments. In this respect, hydrogen

should be considered as a major constituent for these films rather than

an "additive" . This is a very important area for detailed studies.

3 . 2 Substrate Effects

The difference in the oxidation behavior of polycrystall ine Ta

foil and vacuum deposited Ta film on silicon, described above, demonstrates that the substrate influences the structure of the oxide film.

In this particular case, the Ta film was also polycrystalline but the

crystallite size was very small (^100A) and the film was porous. Even

so, vitreous oxide films formed. A substrate effect was observed even

with Si0 2 films: oxide films grown on sandblasted single crystal Si

surface had crystallites, whereas the film grown on smooth surfaces

were vitreous. 11 The most significant substrate effect is the

tremendously increased corrosion resistance (by 5-6 orders of

magnitude) of amorphous metals relative to polycrystal 1 ine metals of

identical composition to be discussed below. Substrate effects are probably involved in the enhanced corrosion resistance of pure iron

which was implanted with Ar ions. 29 This was attributed to the

thickening of the air-formed oxide film as inferred from electrometric

reduction. However, it is possible that the structurally disturbed or even amorphous metal surface, resulting from the ion bombardment,

behaves somewhat similarly to amorphous metals. As a result, a noncrystalline passivating film may form with greater ease. It is

expected that, even when the main propose of ion implantation is to

introduce an additive which improves the passivation behavior, some effects due to surface disordering in the metal structure will arise.

Crystal nucleation rate in the noncrystal 1 ine oxide film and,

hence, the stability of the noncrystalline structure depend very much

on the substrate. This indicates that the formation and stability of

a noncrystalline film are determined by the balance between the com- peting processes of its growth and crystall ization. '

3.3 Thin Film Effects

An important, but frequently overlooked, point is the implications of the fact that passivating films are very thin. The usual

26
considerations derived from bulk behavior are not necessarily applicable

to thin film phenomena. For example, PtSi can be formed at ^200°C by

an interface reaction between Si substrate and thin (several hundred

Angstroms) Pt film whereas, according to the phase diagram, the minimum

temperature should be above 700°C. 30 Also, an oxidizing reaction

due to oxygen diffusion through the Ta 2 0 5 takes place at the

Si/Ta 2 0 5 interface at a temperature which is not sufficiently

high to oxidize bare silicon. 15 Kinetic rather than thermodynamic

considerations are crucial in thin film phenomena where the

interface energy is often high and overcomes other energy

barriers

.

These "thin film effects" further complicate the situation in

noncrystall i ne oxide films which is already complicated due to the

lack of crystallinity. The structural information obtained from

x-ray or electron diffraction is not enough to characterize these

films. There is a great need for using techniques which give information on short-range order and on the nature of the chemical bonds.

These studies should complement electrochemical measurements. The

complexity arising from the thinness of passivating films and from

their noncrystall i ne structure can be unraveled only by using

sophisticated techniques and, just as importantly, sophisticated

and unconventional concepts and ideas.

4. NONCRYSTALLINITY AND PASSIVITY

There are various indications that the special properties of

noncrystalline films described above affect the formation stability

and repai rabi 1 i ty of passive films. It will be the task of this

section to examine the evidence that noncrystal 1 ine films provide

more "perfect" passive films.

4 . 1 Formation of More Protective Passive F ilms

Both the rate of growth and the limiting thickness of passive

films should be affected by the degree of crystallinity. This is so because these factors are dependent on the ability of anions or cations

to move through the growing film. As pointed out earlier, this

retardation in ionic mobility occurs because noncrystal 1 ine films have

few defects or grain boundaries that enhance ionic movement. There

are a number of experiments that can be interpreted as supporting

this concept.

First, consider a situation where an alloy forms a noncrystalline

film. An example is the Fe~Cr alloys shown by McBee and Kruger 2 to

form noncrystalline films on more and more surface grain orientations

as the Cr content goes up. In Table I is shown a comparison of the

limiting thicknesses formed in the passive potential region on alloys

containing different Cr contents. It can be seen that with increasing

Cr and the increasing noncrystal 1 ini ty of the passive film, the limiting

film thickness is decreasing. Moreover, as the thickness of the passive

film that forms on an alloy surface is decreased, its ability to protect

that surface increases. High chromium alloys are, of course, more corrosion

resistant than are low chromium ones.

27
As was mentioned above, amorphous alloys exhibit remarkable

corrosion resistant behavior. 31 For example, Hashimoto et al. 31 have

shown that the passive current for an amorphous Fe-Cr-P-C alloy is over

two orders of magnitude lower than that for the same alloy that has

been converted from a metal glass to the usual crystallized metal.

Such superior ability to lower the passive current density must reside

in the superior ability of the film formed on the metal glass to

decrease the flow of ions.

As discussed earlier above, hydrogen appears to play a beneficial

role in affecting the properties of noncrystalline films or their

ability to form. For example, Ti forms a much more protective film

in wet oxygen as compared to the film formed in dry oxygen. 4 A number

of workers have provided evidence for hydrogen or "bound water"

in both crystalline and noncrystalline passive films. 3

, 32 For example the

passivity associated with the outer layers of the passive film on iron

was found to depend on the role of hydrogen in promoting the formation

of the crystalline y-Fe 2 03. 33 Therefore, the role of hydrogen may not

be only related to the promotion of noncrystall i ni ty or its effect on

the structure of noncrystall ine films.

4.2 Resistance to Breakdown

Just as noncrystallinity appears to have some beneficial effect

on the formation of a more protective film in the passive region of

the anodic polarization curve, there are also indications that when

a metal is at a potential above the critical potential for pitting, the

ability to resist breakdown leading to pitting or crevice attack is

enhanced by the existence of a noncrystalline film on its surface.

Fig. 1 shows a comparison between the breakdown behavior on Fe (crystalline film) and on Fe-Cr (noncrystalline film). The superiority of

the noncrystalline film with regard to resistance to breakdown is

reflected in its greater time for breakdown even while being thinner

and being exposed to higher concentrations of damaging species (chloride

ions). Also of great interest is the finding that for the crystalline

passive films on Fe, one can remove Cl ions from the film by replacing

the chloride containing solution with one that does not contain chloride,

whereas for the noncrystalline films on Fe-Cr, the Cl cannot be removed

in a similar fashion. Therefore, noncrystalline films resist breakdown

by making entry by damaging species more difficult, but they also inhibit

exit once some penetration has taken place. This would imply that

whereas entry or exit of Cl - occurs via defects or grain boundaries in

crystalline films, it may only enter the noncrystalline films by becoming

part of their "glass-like polymeric" structure. Such stable structures,

as the results in Fig. 1 show, are both difficult to enter and difficult

to exit.

Similarly, the breakdown of a film in a crevice is affected by the

crystal 1 ini ty of the passive film. Fig. 2 shows results using an optical

technique described elsewhere. 36 It shows that as one increases the Cr

content, the tendency to breakdown as reflected by the change in the

28
optical parameter, a, (relative phase retardation) is retarded. Ambrose

and Kruger 36 have attributed this decrease in A to the onset of conditions

in a crevice that lead to breakdov/n.

Two other indications that passive film crystallinity can affect

breakdown are provided by the work of Devine and Wells 37 and Okamoto 3. The former observed that no pitting occurred because of breakdov/n of the

films on amorphous metals, presumably (our interpretation) because these

films were noncrystal 1 ine. Once the metals were crystallized by heat

treatment, however, pitting took place, presumably because the films

formed on the crystalline substrate were no longer crystalline. Okamoto

showed that the induction time for pitting of an austenitic stainless

steel could be decreased if its noncrystall ine film were annealed.

This behavior is exactly the opposite to that found for iron by Ambrose

and Kruger 38 who found that annealing increased the time to breakdown

(induction time for pitting). A plausible explanation for these opposite

behaviors is that annealing promotes crystal! ization of the noncrystalline

film on the austenitic stainless steel, making breakdown easier by providing more paths for Cl ingress. For the crystalline film on Fe, hov/ever,

annealing removes defects such as cation vacancies and thereby makes

ingress of Cl~ and thus breakdown more difficult.

4.3 Duct i 1 ity

The last section was concerned with chemical breakdown. There is

also the possibility of mechanical breakdown the rupture of passive

films by applied stress. This form of breakdown is of great importance

for stress corrosion cracking of metals. Noncrystal 1 ini ty of passive

films can play an important role affecting this kind of breakdov/n through

its influence on the ductility of the film. A more ductile film would

have a better chance for avoiding mechanical breakdov/n. Vermilyea 39 has

stated, for example, that stress corrosion cracking is not possible if

the passive film is plastically ductile. It should be pointed out also that some theories of chemical breakdov/n leading to pitting also postulate film

rupture by chemical forces 40

, and thus ductility may be important for

all forms of breakdown.

There are some indications in the literature that degree of

crystallinity does affect film ductility. Bubar and Vermilyea 41 showed that

anodic Ta 2 0 5 films lost their ductility after crystallization. Bubar

and Vermilyea 42 also have shown that there is a wide variation in the

ductility of passive films formed on a number of metals. Films formed

on Ta, which is known to form noncrystal 1 i ne films, exhibited deformations of as much as 50% before fracture. Also, the passive film on

304 stainless steel (noncrystall ine) exhibited better ductility than that

on crystalline iron. Aluminum, which may or may not form noncrystal 1 i ne films, formed films less ductile than those on Fe. Leach and Neufeld 43

, however, found easy deformation of films on A1 under an applied field.

It is thus not certain that the ductility of passive films can only be

attributed to noncrystall i ni ty. Other factors such as the presence

of hydrogen in the film may also play a complicating role.

29
4.4 Repassivation

When breakdown occurs, repassivation rate can play an important

role in determininq whether passivity can be restored or whether

localized attack such as cracking, pitting or crevice corrosion can

be initiated. There are no direct experiments to our knowledge that

indicate that noncrystal 1 i ne films reform at a greater rate than

crystalline ones do, but one can obtain indications of this if one compares

the repassivation of low carbon steel (crystalline film) to austenitic

stainless steel (noncrystalline film). Using the technique of triboel 1 i psometry 44

, 45, Ambrose and Kruger have obtained repassivation ratios,

Rp, for these metals. Rp is a measure of the efficiency of repassivation;

it is the ratio of the total current during repassivation (the current

consists of both dissolution and film growth components) to the thickness

of the reformed passive film for a given instant of time during the

repassivation event. Fig. 3 shows that the time of repassivation and

the magnitude of Rp are considerably lower for the stainless steel

than they are for low carbon steel. These results thus indicate that

the rate of passive film regrowth and the effectiveness of retarding

metal dissolution are enhanced for alloy surfaces that form non- crystalline films.

5. CONCLUSIONS

Owing to the lack of grain boundaries, noncrystalline oxide

films have better passivating properties than polycrystal 1 ine ones.

Bond and/or structural flexibility is an important factor in determining the ability to form a noncrystall ine structure and its

stability: This point was demonstrated by the example of vitreous

Si0 2 and Ta 2 0 5 as well as by ligand formation of transition metal

ions. The flexibility of the noncrystalline structure ensures a good

accommodation at the oxide/substrate interface without requiring an epitaxial relationship. Such an interface is less prone to chemical

attack than an imperfect one. High degree of short range order in

vitreous oxide films also contribute to their chemical resistance.

The flexibility of the structure is responsible for the increased

ductility of noncrystall i ne passivating films which is crucial from

the viewpoint of pitting and stress corrosion.

The structural /bond flexibility and, hence, the tendency toward

noncrystall inity can be increased by additives; the effect of Cr in

the passivation behavior of stainless steel is an example. In this

respect, hydrogen also plays a great role by forming M-OH groups,

hydrogen bonds, and H 2 0 ligands.

The formation and stability of noncrystalline oxide films is

a kinetic rather than thermodynamic problem as demonstrated, for

instance, by the influence of substrate. Due to the increased

stability of noncrystalline passivating film, amorphous alloys

exhibit orders of magnitude greater corrosion resistance than the

polycrystal 1 i ne alloys of identical composition.

30
Thermodynamic and other concepts used in connection with bulk

crystalline solids are not a priori applicable to thin and non- crystalline passivating films. Many aspects of thin noncrystal 1 i ne films

are determined by unique interface reactions and unique properties of a noncrystalline structure.

There is evidence in the literature that noncrystal 1 ine passive

films are 1) more protective, forming films which exhibit lower limiting

thicknesses and lower passive current densities; 2) more ductile, 3)

less susceptible to breakdown by damaging species and 4) capable of

reforming at greater repassivation rates than crystalline films.

ACKNOWLEDGEMENTS

The contribution to this paper of A.G. Revesz is based upon work

performed in COMSAT Laboratories under the sponsorship of the Communications Satellite Corporation. The contribution of J. Kruger was supported by the Office of Naval Research under contract NAOMR 18-69

NR036-082

.

Thanks are due to F. Fehlner for helpful comments on the manuscript.

31
References

1. T.P. Hoar, J. Electrochem. Soc. 117 , Mo. 1, p. 17C (1970).

2. C.L. McBee and J. Kruger, Electrochem. Acta, 1_7, 1337 (1 972).

3. G. Okamoto, Corrosion Science ]J3, 471 (1973).

4. J.S.L. Leach, Surface Science 53^ 257 (1975).

5. F.P. Fehlner and N.F. Mott, Oxidation of Metals 2_, 59 (1970).

6. W.H. Zachariasen, J.Am. Chem. Soc. 54^ 3841 (1932).

7. K.R. Lawless, Repts. Progr. in Physics 37, 231 (1974).

8. M.J. Pryor, Oxidation of Metals .3, 271 (1971 ).

9. L. Young, Anodic Oxide Films , Academic Press, New York (1961).

10. P.J. Harrop and D.S. Campbell, Thin Solid Films 2_, 273 (1968).

11. A.G. Revesz, Phys. Stat. Soc. 24_, 1 15 (1 967).

12. M.M. Atalla, E. Tannenbaum, and E.J. Scheibner, Bell System

Technical J. 30, 749 (1959).

13. A.G. Revesz, J. Mon-Cryst. Solids 1_1_, 309 (1973).

14. A.G. Revesz, J. F. Allison and J.H. Reynolds, COMSAT

Technical Review, 6., 57 (1976).

15. A.G. Revesz and T. Kirkendall, J. Electrochem. Soc. 1 23 , 1514-1519.

16. F.P. Fehlner, J. Electrochem. Soc. 119 , 1723 (1972).

17. R.L. Myuller, in Solid State Chemistry (O.U. Borisova, ed.),

translated from Russian, p. 1, Consultants Bureau, New York,

(1966).

18. C.S.G. Phillips and R.J.P. Williams, Inorganic Chemistry , Vol

.

1, p. 480 and Vol. 2, pp. 189-239, Oxford University Press,

New York, (1965).

19. P. Kofstad, High Temperature Oxidation of Metals , J. Wiley,

New York (1966).

20. N.C. Stephenson and R.S. Roth, Acta. Cryst. B27 , 1037 (1971).

32
21.

R.S. Roth and J.L. Waring, J. Res. MBS 74A, 485 (1970).

22. J.S. Anderson, J. Chem. Soc. (Dalton), 1973, p. 1107.

23. B.G. Hyde, A.N. Bagshaw, S. Anderson, and M.O. Keeffe, in

Annual Review of Materials Science , Vol. 4 (R.A. Huggins

et al , eds . ) ; p. 43, 1 974.

24. A.F. Wells, Structural Inorganic Chemistry , p. 459 Clarendon

Press, Oxford, 1975.

25. V. Ashworth, W.A. Grant, and R.P.M. Proctor, Corrosion

Science 1_6, 661 (1976).

26. A.F. Wells, op.cit . , pp. 526-529.

27. C.S.G. Phillips and R.J.P. Williams, op. cit . , Vol. 1, p. 533.

28. A.G. Revesz, J. Electrochem. Soc., in press.

29. V. Ashworth, W.A. Grant, R.P.M. Proctor, and C. Wellington,

Corrosion Science, 1_6, 393 (1976).

30. W. Mayer and K.N. Tu, J.. Vac. Sci. Tech. Vol. 1J_, 1074.

31. K. Hashimoto, K. Osada, T. Masumoto, and S. Shimodaira,

Corrosion Science 16, 71 (1976).

32. H.T. Volken, J. Kruger and J.P. Calvert, Corrosion Science,

8, 103 (1968).

33. C.L. Foley, J. Kruqer, and C.J. Bechtoldt, J. Electrochem.

Soc. m, 944 (1967).

34. C.L. McBee and J. Kruger, Nature Physical Sci. 230 , 194 (1971).

35. C.L. McBee and J. Kruger, Proceedings Japan-U.S.A. Symposium

on Passivity and Breakdown of Passivity of Iron Base Alloys,

R.W. Staehle and H. Okada, eds.. National Association of

Corrosion Engineers, Houston, 1976, p. 131.

36. J. Kruger and J.R. Ambrose, Nat. Bureau of Standards

Report NBSIR 74-583, Sept. 1974.

37. T.M. Devine and L. Wells, Scripta Metallurgia 10, 309 (1976).

38. J.R. Ambrose and J. Kruger, Proc. 4th Int. Cong, on Metallic

Corrosion, National Association of Corrosion Engineers,

Houston, 1972, p. 698.

39. D.A. Vermilyea, Physics Today, 29, 23 (1976).

40. J. Kruger, see Ref. 35, p. 91.

33
41 . S. F. Bubar and D.A. Vermilyea, J. Electrochem. Soc. 114, 882 (1967)

42. S.F. Bubar and D.A.

892 (1966).

Vermilyea, J. Electrochem. Soc. 113,

43. J.S.L. Leach and P. Neufeld, Corrosion Sci. £, 225 (1969).

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

45. J.R. Ambrose and J. Corrosion, NationalKruger, Proc. Fifth Int. Cong, on Metallic

Association of Corrosion Engineers,

Houston, 1974, p. 406

34
Figure Captions

Figure 1. (a) Spectro-el 1 ipsometric measurements of changes in A,

relative phase retardation, and Y, relative amplitude

reduction, for iron after introduction of 0.005N chloride

(t=0), and removal of chloride from the solution. A, y, and current are seen to return to normal passive values once chloride ion is removed. 410 nm wavelength light

was used. From (34).

(b) Spectro-el 1 ipsometric measurements of changes in A and t for Fe-20 Cr after introduction of 0.08N chloride

(t=0), and removal of chloride from solution, a and f

do not return to passive values once chloride ions are removed. 450 nm wavelength light was used. From (35).

Figure 2. Changes in relative phase retardation, 6A, as a function

of time measured within a crevice, using a glass plate

to produce the crevice for Fe-14%Cr and Fe-17%Cr alloys

exposed to IN NaCl solution. From (36).

Figure 3. Comparison between the variation in Rp (repassivation

ratio determined by tri boel 1 ipsometry ) in 1 N NaCl for

304 stainless steel (From 36) and low carbon steel (From

45).
TABLE I - Variation in Limiting Film Thickness As A

Function of Chromium Content (from Ambrose and Kruger,

unpublished) In Iron Chromium Alloys and Crystallinity

of The Passive Films Formed On These Alloys (from 2).

I Cr Degree of Limiting

Crystallinity Film Thickness

of Passive Fi 1m (&)

0 Well oriented 36

spinel

5 Well oriented 27

spinel

12 Poorly oriented 21

spinel

19 Mainly non- 19

crystal 1 ine

24 Completely non- 18

crystalline

36
(a)

</>

i/> a>o>

a> Q

86.5 -

86.0 -

w 24 V)at

u.D>

a> O^ 23 -

am ^ _ o ° ° °°oo o oo

000

10 . too

a.

Chloride

Added

o o o

Chloride

Removed

o o o

o

OOOOOO O O

0 0

X = 45008

o o

ooo

oo Fe-20 Cr

1

0 40 80 120

TIME (min.)

160 200

(b)

37
T~

r Z

38
rj

3

3

o

o

c

3D

“O

r-' 2

1

39
PART III To be submitted to Corrosion J.

Effect of Chromium Depletion on the Repassivation Kinetics of

18-10 Austenitic Stainless Steel in Sodium Chloride Solutions

John R. Ambrose

National Bureau of Standards

Institute for Materials Research

Washington, D. C. 20234

A great deal of effort has been directed toward obtaining an understanding of the relationship that appears to exist between the susceptibility

of AISI 304 austenitic stainless steel for intergranul ar stress corrosion

cracking (ISCC) and the tendency for this particular alloy to become

sensitized as a result of welding operations. The importance of this re- lationship has been demonstrated upon careful analysis of the phenomonology

of failures of AISI 304 pipe material exposed to high purity water in

BWR nuclear reactor systems (1). In these systems, failure occurred either

in the heat affected zones (HAZ) of weldments or in severely furnace sensitized

material

.

Although it is but one of several suggested mechanisms, the idea that

the occurrence of carbides in grain boundaries as a result of improper heat

treatment can be related to the phenomenon of sensitization is most interesting.

The formation of these metallic carbides is presumably associated with a correspond!' ng depletion in chromium in areas immediately adjacent to the

grain boundaries where the particles have formed. Depletion of chromium to

concentrations below that required to sustain passivity would account for

many of the properties which have been displayed by material which has been

sensi ti zed

.

However, the relationship between ISCC and sensitization becomes more complex if it is considered that sensitization appears to be a necessary

but not sufficient criteria for controlling susceptibility of AISI 304

stainless steel to ISCC. This conclusion arises from the fact that out of some 17,000 welds, less than 100 have experienced any failure. Thus,

techniques which measure only degree of sensitization will not necessarily

yield information by which susceptibility to ISCC can be ascertained

(i.e. ASTM Standard Procedure A-262, practices A, C, and E). Another

parameter which has been shown to be related to the susceptibility of a material to stress corrosion cracking is the rate at which a bare metal

surface will repassivate in the environment to which it has been exposed.

The rationale behind this relationship of SCC to repassivation kinetics

is incorporated into the mechanism described in other publications (2, 3).

It has been known for some time that overall corrosion resistance of

stainless steels is related to the concentration of chromium in the alloy

and that a minimum concentration of chromium (approximately 12%) is required

to maintain a passive condition. It also has been shown that a relationship

does exist between repassivation kinetics and susceptibility to SCC for AISI

304 austenitic stainless steels (4). What is not known, however, is what

relationship exists between 1) chromium depletion of austenitic stainless

steels and repassivation kinetics in solutions known to cause SCC and between

- 40 -
2) sensitization and repassivation kinetics in these environments. This

paper describes the results of a study into the effect of chromium depletion

on the repassivation kinetics of Fe-18Cr-10Ni austenitic stainless steels.

Experimental

As described in a previous report (5), a series of ternary Fe-lONi-xCr

alloys were prepared where x was varied between 3 and 18 weight percent in

increments of 3%. Although considerable effort was expended to retain the

austenite structure in these alloys, the final structure in all specimens

other than 18%Cr was a martensite-like phase dispersed in a ferritic steel

matrix. There is, of course, the question as to what effect the alteration

of structure would have on repassivation kinetics of a particular alloy

composition but this will have to be determined at a future time. These

studies should, however, provide insight into how alloy composition is

related to repassivation kinetics.

Cylindrical segments were cut from the prepared rod material and

mounted in a Buehler expoxy mount. The finished specimens were drilled

and trapped for electrode assembly. Repassivation kinetics were determined using a modification of the tri bo-el lipsometric technique in which

the anodic current decay was recorded as a function of time on an oscilloscope. All experiments were performed in a 1.0N NaCl solution in which

no attempt was made to deaerate the system. Measurements were made for

each alloy at its respective corrosion potential and at several other

more noble potentials, including values above the critical pitting

potential of that particular alloy.

The experimental procedure in each instance was the same: After

electrode installation in the electrochemical cell and connection of

electronic leads, solution was added to the cell and the rest potential

allowed to stabilize. The tribo-el 1 ipsometri c abrasion apparatus was activated and the metal surface abraded for approximately 5 seconds following which the abrasion wheel was retracted and repassivation was monitored

by recording the current decay transient on an oscilloscope. The process

was repeated for a given alloy potentiostatted at several more noble

potentials. Solution was changed for each alloy and the experimental

procedure repeated.

Results and Discussion

In figure 1 are plotted current decay against time for the Fe-18CrlONi alloy at several potentials. Current decay kinetics were virtually

independent of applied potential. In Figure 2 are plotted results from

repassivation kinetics of Fe-15Cr-10Ni alloy. Initial current decay

kinetics are independent of potential, but above the critical pitting

potential, at 249 mV SHE, current decay abruptly slows as onset of pit

growth begins. Results are not shown for the Fe-1 2Cr-l ONi alloy since

the measured currents at the corrosion potential were inconsistent with

results obtained from other alloys. Microscopic inspection of the specimen surface revealed crevice corrosion attack at the interface between

the metal and the expoxy mounting resin due probably to an imperfect seal

of mount to metal. Results of experiments in which this alloy has been

tested will be presented in future work.

- 41 -
In Figure 3 are plotted the current decay kinetics for the 9% chromium

alloy. In this case, there is a marked difference in the potential dependence of the repassivation rate. This effect is particularly pronounced as the potential is increased from 150 to 250mV SHE. Above 250mV, the critical

potential for pitting. is exceeded. At 350mV, after about 45 milliseconds

there is an abrupt increase in the rate of current decay. This effect is

probably due to formation of corrosion product at the solution interface

which restricts anodic dissolution.

Figure 4 shows the repassivation kinetics for the 6% alloy. Current

decay kinetics for this and the 3% alloy (figure 5) are markedly decreased

Again, as for the 91 alloy, there is a potential dependence exhibited for

the rate of current decay. In addition, at the higher applied potential,

there is an inflection in the current decay plot at 165 mV at about 32 ms.

A similiar increase in the rate of current decay is noted for the 3 % alloy

at 9mV and 25ms. It is interesting to note that the point in time at

which these inflections occurred in the 9, 6, and 3% alloys corresponds

to approximately the same quantity of consumed change for each alloy, ^0.5

coulombs. This would further support the view that repassivation is due

in part to the formation some insoluble substance, which precipitates

at the metal surface when the solubility product of that substance has

been exceeded.

In Figure 6 are compared repassivation rates for the 18, 15, 9, 6

and 3% alloys at their respective open circuit or corrosion potentials.

It is immediately obvious that repassivation is exceptionally difficult

for the 6 and 31 alloys. The difference in repassivation rates becomes

much more pronounced under anodic polarization. In Figure 7 are compared current decay kinetics for all alloys of the series at a given

applied potential of 142mV SHE. From this plot, the set of alloys can be

divided into three groups based upon their rates of repassivation: group

I - 18, 15, 12% Cr, fastest repassivation kinetics; group II - 6, 3% Cr,

slowest repassivation kinetics; and group III - 9% Cr; intermediate repassivation rates.

From these results it appears likely that the rate of repassivation

of this family of stainless steels is proportional to the concentration

of chromium in the alloy. It is particularly worthwhile to note that an area of intermediate repassivation kinetics does exist, between 12 and 6%

chromium. This is interesting because such an intermediate value of re- passivation rate has long been suggested as being a criteria for propagation of sharp cracks during ISCC.

The fact that an intermediate region of passivation behavior exists

also explains why sensitization is considered a necessary but not sufficient

condition for susceptibility to ISCC. Areas of the heat affected zone must

be of a certain chromium concentration (degree of sensitization) and in an area of the metal where metallurgical factors are such that crack advance

is favored.

- 42 -
Of secondary interest is the observation that, under conditions of

restricted repassivation by the original passive film, current decay is

apparently affected by the formation of another substance at the metal

surface. Perhaps as has been previously suggested, this substance is a salt film of limited solubility under these conditions of repassivation.

Two considerations should be made in treating the results obtained

from this study to date:

1) The technique does allow detection of intermediate values of

repassivation behavior. Techniques based upon an analysis of time to

failure or examination of a current-potential plot will not permit such

detection since the measured parameters are based upon a cumulative series

of events. No one step of the process can be isolated from the rest.

By using a technique based on repassivation as being the most important

step in controlling ISCC susceptibility, measurement of a parameter

which is directly related to the repassivation kinetics of exposed

metal will provide a more meaningful diagnostic tool.

2) Although the technique can be used to detect values of the

chromium concentration where an alloy becomes susceptible to ISCC, it

is not known at this time if differences in repassivation kinetics can

be measured when the chromium depletion is restricted to minute areas

at the surface of the specimen - i.e. adjacent to the grain boundaries

of a polycrystal 1 i ne material. Specimens have been produced in which the

material has been heat treated to several degrees of sensitization.

Results of their measured repassivation kinetics should answer the

question.

Conclusion

1) Repassivation kinetics of Fe-Cr-Ni stainless steels are affected

by variations in the chromium concentration of the alloy.

2) There exists an intermediate range of repassivation behavior

between 12 and 6% chromium concentration which may be tied to the question

of how sensitization and susceptibility to ISCC are related.

3) Even though chromium depletion is related to repassivation kinetics,

it is not known whether this technique can detect variations in the repassivation kinetics of material which has been heat treated to different degrees

of sensitization. These results will be the subject of the next report.

- 43 -
References

1. NED6-21000, Report prepared for Electric Power Research Institue

by A.J. Giannuzzi, General Electric Project Engineer, BWR Reactor

Systems Department.

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

3. J. R. Ambrose and J. Kruger, J. Electrochem. Soc. 121 , Mo. 5., 599

(1974).

4. Office of Naval Research, NAONR 18-69 NRS Report No.

NBSIR 74-583, September 1 Q 7*.

5. Ibid., NBS Report No. NBSIR 76-1170, November 1976.

44
Figure Captions

Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Figure 6

Figure 7

Plot of anodic current in milliamperes

seconds for Fe-18Cr-10Ni in 1.0 N NaCl

518 mV SHE.

Plot of anodic current in milliamperes

seconds for Fe-15Cr-10Ni in 1.0 N NaCl

and 249 mV.

Plot of anodic current in milliamperes

seconds for Fe-9Cr-10Ni in 1.0 N NaCl

and 350 mV.

versus ti me in mi llisol ution at HE!, 218,

versus ti me in mi 11 i - solution at -51 ) 49,

versus ti me in mi Hisolution at 150, 250,

Plot of anodic current in milliamperes versus time in milliseconds for Fe-6Cr-10Ni in 1.0 N NaCl solution at -35, 65,

165 mV SHE.

Plot of anodic current in milliamperes versus time in milliseconds for Fe-3Cr-10Ni in 1.0 N NaCl solution at -91 and

9 mV SHE.

Comparison of repassivation kinetics for series of Fe-XCr-lONi

(x= 18, 15, 9, 6, and 3%) at their respective open circuit

potentials.

Comparison of repassivation kinetics for series of Fe-XCr-lONi

(x= 18, 15, 12, 9, 6, 3%) at 142 mV SHE.

45
Figure 1

46 -
.

Figure 2

- 47 -
Figure 3

48

Figure 5

- 50 -
Figure 6
Figure 7

- 52 -
PART IV To be published in Proceedings of 4th International

Symposium on Passivity

Composition of Surface Films Formed During the Repassivation of Iron and IronMolybdenum A1 Toys

J. R. Ambrose

Institute for Materials Research

Metallurgy Division

National Bureau of Standards

Washington, D.C. 20234

Molybdenum has been in use for some time as an inhibitor of localized

corrosion, both as an alloying addition and as a salt dissolved in the solution

to which the metal is to be exposed. Although the phenomonology of the inhibition has been characterized, little progress has been made in developing

a comprehensive understanding of the inhibition mechanism.

Recent spectroscopic surface profile analyses require new explanations for the influence of molybdenum on passivity and how this passivity

is related to properties of surface films. Auger Electron Spectroscopy (AES)

data have revealed an apparent surface depletion of molybdenum, which

occurred in conjunction with an increased resistance to localized corrosion

attack (1). One possible explanation has been developed from results obtained

from a study into the effect of molybdate ion additions on pitting of iron

in chloride solutions (2). Inhibition of propagation of pits appeared to

be related to the accumulation, in repassivated pits, of molybdenum, which

was thought to be in the form of an iron molybdate salt film. If formation

of such a film were responsible for the resistance of iron to pit

propagation in the presence of molybdates, then it is under those con- ditions of restricted mass transport that are most likely to exist during

propagation of localized corrosion attack that such films would be expected

to be the most stable and the most protective. In order to study repassivation under such conditions, a crevice electrode system with a window

was conceived.

This system was designed so that simultaneous measurement of optical

variations and electrochemical parameters could be made during the onset

of crevice corrosion attack or during cathodic depassivation - anodic re- passivation cycles within the crevice. However, location of the various

electrodes-working , reference, and auxiliary - was such that current

distribution has never been uniform over the exposed surface of the working

electrode. Thus, until now, results have been limited to analyses of

current decay curves obtained during repassivation transients. It was possible, though, even without obtaining measurement of film growth kinetics,

to evaluate the contribution of alloyed molybdenum to the repassivation

kinetics of iron in chloride solutions under the influence of a crevice

and to speculate on the mechanistic implications.

Molybdenum affects the repassivation behavior of the alloys in

two ways

:

a) the initial rate of current decay is proportional to the

concentration of molybdenum in the alloy and

-53 -
b) the length of time during which the current continues to

decrease is also proportional to the concentration of molybdenum in the alloy. Longer time intervals lead to lower

values of crevice corrosion current during exposure.

Based on these electrochemical measurements, it was suggested that the

stability of a molybdate salt film was a contributing factor to the

repassivation kinetics of i ron-molybdenum alloys and thus directly re- sponsible for the observed improvement in corrosion resistance upon

introduction of molybdenum into ferrous based alloys (4).

Critical questions remain - What is the nature of the repassivating

film under conditions where crevice corrosion occurs, under what conditions

does it form, and how fast does it grow? In order to answer the first

question the crevice electrode was redesigned so that more uniform current

distribution could be achieved. Following repassivation in an environment

known to produce crevice corrosion, specimens were removed and the surfaces were analyzed by x-ray photoelectron spectroscopy (XPS) to determine

film composition relative to the concentration of molybdenum in the alloy.

An AEI ES 200 XPS spectrometer was used for these studies; analysis was performed by Nils E. Erikson of the Surface Science Division, National

Bureau of Standards. The results of this study, when coupled with re- sults presented in preceeding reports (2, 3), give some insights into the

role of molybdenum as an inhibitor of corrosion.

Experimental

The crevice electrode assembly was virtually identical to that

already described (3) except that the glass disc used to produce the crevice

was coated on one side with a film of tin oxide. Silver epoxy was used

attach a piece of platinum wire for electrical connection. The electrically conducting tin oxide served as an auxi 1 iary electrode located within the

crevice. This particular orientation of counter to working electrode,

coupled with the position of the reference electrode within the crevice,

permitted a better current distribution over the working electrode surface

than has been achieved to date.

Specimens of Fe-lMo and Fe-5Mo, machined from the same material

described in previous reports (2, 3) to 2.5 cm x 0.15 cm x 0.62 cm, were

so designed that they could be placed in the chamber of an XPS spectrometer. A section of teflon rod was slotted in one face to hold the specimens, then drilled and tapped to receive necessary electronic connections.

Surfaces were polished through SiC papers to a 1 micron diamond paste

finish. Spectra were obtained for both compositions before exposure to

any solution; XPS scans were performed to determine composition at the

surface of the air-formed film and again following argon ion sputtering

to remove oxide films.

Specimens were then assembled in the crevice electrode system and

exposed to an 8.41 pH borate buffer solution. Cathodic reduction at -758

mV SHE preceded repassivation at a potential of 442 mV SHE. Quantity of

-54 -
repassivation was monitored coulometrical ly and, when 1.15 x 10~2 coulombs

had been consumed, the potentiostat was turned off, the electrode disassembled as rapidly as possible, and the working electrode rinsed with spectro

grade methanol and dried in a stream of compressed air. The specimens were

removed from the teflon holder, wrapped in a piece of lens tissue, placed

in a coin envelope, and transferred to the laboratory where the XPS analysis

was performed.

Following analysis, specimens were repositioned in the teflon

mounting adapter and returned to the electrochemical cell. This time,

the solution used contained 1 O' 2 M sodium chloride in the 8.41 pH borate

buffer. The experimental procedure was identical to that described above.

However, there was some over run in the charge consumed due to the decreased

repassivation kinetics in chloride containing solution. The consistancy

of the consumed charge for all experiments was within ±2%. Following

solution exposure, XPS analysis was repeated.

Results

Compositions of Film on Fe-Mo Alloys

An XPS spectrum obtained from an air-formed film on molybdenum is

shown in figure 1. By comparing Mo 3d peak positions in this spectrum

with those in a spectrum obtained following film removal by argon ion

(Ar+

) sputtering, it is possible to distinguish between elemental and

combined molybdenum. As a first approximation, quantitative relationships

were derived from calculation of peak height ratios of the respective

states. For the two Fe-Mo alloys, quantitative results are presented in

figure 2 as ratios of elemental iron and molybdenum to their respective

combined states. Also presented in figure 2 are the ratios of total molybdenum to total iron. Mo/Fe peak height ratios for the oxide free surfaces

(Ar+

) of the Fe-5Mo and Fe-lMo alloys are themselves in a ratio of 5:1

(3.0 to 0.6) which indicates that the peak height measurement is appropriate for estimating surface composition ratios. In fact, the Mo/Fe ratios

for the two alloys were in a 5:1 ratio for all surface conditions studied - in figure 2 these are denoted as:

Ar

+

- film free surface following Argon ion sputtering

Air film - air formed film following polishing

No (Cl) - anodic film grown in absence of chloride ion

(Cl) - anodic film grown in presence of. chloride ion.

However, Mo/Fe ratios for the respective air formed films on each alloy

were substantially less than in the substrate. As figure 2 shows, (Mo/Fe)

5Mo

= 2.2, and (Mo/Fe)]^0

= 0.45. This ratio decrease indicates a surface

depletion in the amount of molybdenum relative to iron in the film.

This apparent molybdenum depletion is even more pronounced in

anodic films grown in the absence of chloride iron. Mo/Fe ratios for

Fe-5Mo and Fe-lMo were 1.4 and 0.3, respecti vely . Furthermore, from figure

- 55 -
2, comparison of peak height ratios for elemental to combined states for both

iron and molybdenum indicate that the anodic film is thicker than the air

formed film.

The last set of results are from anodic polarization experiments

in solutions containing 10 2 M NaCl . All ratios, Fe/ m )/Fe(ox) , Mo( m

)/Mo(

ox )

, and

Mo/Fe, have increased in value over those from experiments in which no chloride was present in the solution. The ratios all approach those values

determined for the respective air-formed films. It should be remembered

that in these chloride containing solutions, repassivation rates for each

alloy were substantially less than in non-chloride environments.

Bonding Character!' sties of Molybdenum in Films on Fe-Mo Alloys

A comparison of the spectra for air-formed film to the anodic

film on Fe-5Mo obtained following repassivation without chloride ion

is given in figure 3. Notice the appearance of a peak intermediate

between the two higher kinetic energy peaks as well as a disappearance

of a low kinetic energy peak at the left side of the spectrum.

Since results shown in figure 2 indicate a large molybdenum depletion in the anodic film on these alloys, the spectrum obtained for the

anodic film on Fe-5Mo grown without chloride ion could reflect the depletion of high oxidation state (+6) molybdenum compounds. However, when the

spectrum obtained for an anodic film grown on Fe-5Mo in 10 2 M Cl is compared

to that obtained for the air-formed film (figure 4), a return to the

characteri sti cs of the air-formed film is noted. As noted in figure 2,

this change in the XPS spectrum for the anodic film is accompanied by

an increase in the amount of molybdenum relative to iron.

Chloride in Films on Fe-Mo Alloys

The chloride absorption peaks for both alloys following repassivation

in chloride solution at 442 mV SHE are compared in figure 5. There is

substantially more chloride present in the Fe-5Mo alloy than in the IMo

alloy. It should also be noted, however, that the ratio of peak heights

for chloride is not in the same ratio as the concentrations of molybdenum - i.e. not 5:1. Rather, the ratio appeared to be about 2:1.

Di scussion

Results obtained indicate a synergistic relationship between molybdenum and chloride ion during the repassivation of iron and i ron-molybdenum

alloys. Previous studies have been shown that molybdenum has little effect

on the repassivation of iron whether in the form of an alloying addition (3) or as an inhibitor in the solution (2) unless chloride ion is present

in the solution. This study has shown that, unless chloride ion is present in solution, molybdenum becomes depleted in anodic films during re- passivation. A previous study (3) also showed that repassivation kinetics

in chloride containing solutions were proportional to the concentration

of molybdenum in the alloy. A final observation from previous studies (2,

3) v/as that a condition of restricted solution mass transport was required

for molybdenum additions to become effective. The role of the restricted

mass transport condition is not only to create a situation requiring a minimum value of concentration of Mo04 =, but also a minimum value

of Cl ion to synergisti cally interact with the inhibitor ion.

-56 -
The fact that higher molybdenum containing alloys, which repassivated

more rapidly, were found to contain higher amounts of chlorine in the anodic

film further points to a synergistic interaction between the two species.

It is difficult to explain the increased repassivation kinetics Fe-5Mo

over Fe-lMo on the basis of an insoluble molybdate salt film since XPS

results do not show any enrichment of Mo, even under restricted mass transport conditions, for anodic films grown without chloride ion. In the

presence of chloride ion, however, anodic films are enriched in molybdenum, as well as in chlorine, at least with respect to films grown without chlorides in solution.

Therefore, although the formation of a molybdate salt film is not

reasonable, the existance of a chloride salt film is. The properties

of such a chloride salt film should become important in controlling repassivation kinetics and in overall resistance to localized corrosion attack.

The role of molybdenum appears to be in stabilizing a chloride salt film

rather than in affecting repassivation of the substrate directly. Sakashita

and Sato (4) have recently suggested that molybdate ion, by adsorption

on a surface film, can alter transport properties of that film. The formation of a chloride salt film is suggested by the results shown in figure 5 in which chloride ion concentration at the surface is proportional to

molybdenum concentration in the alloy.

Other recently published results on the effect of molybdenum additions to ferric stainless steels (5) have also demonstrated that chloride

ion (as well as alloyed chromium) is required for the beneficial effect

of the molybdenum to be exhibited.

On the basis of these results then, it can be concluded that although

there is a definite increase in the rate of repassivation of iron in chloride

containing solutions as a result of molybdenum addition, this beneficial

contribution is not due to the formation of a molybdate salt film, but

rather is due to the stabilization, by molybdate ion, of a chloride salt

film. In all probability, however, once the repassivation transient is

over, the chloride film will eventually hydrolyze to an oxide film. This

film can also be stabilized by molybdate ion adsorption, as has been suggested by Sakashita and Sato (4). This last statement is also consistent

with what has been observed regarding the inhibition of pit propagation on iron in chloride solutions by molybdate ion (2), in which molybdenum, but

little chloride ion, was found within the repassivated pits.

Conclusions

1. Molybdenum is depleted in surface films on iron-molybdenum binary

alloys.

2. Concentration of molybdenum in surface films is measureably

increased when films are grown in the presence of chloride

ion

.

-57
3. Effective repassivation of iron-molybdenum binaries requires

some kind of synergistic interplay between the molybdate and

chloride ions.

4. It is suggested that Mo04 acts as an inhibitor in two stages:

a) State I - Synergistic interaction with Cl" ion to stabilize

a Cl” salt film.

b) Stage II - Synergistic interaction with an iron oxide to

stabilize a passive film.

- 58 -
References

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

2. Office of Naval Research, NAONR 18-69, "The Role of Passive Film

Growth Kinetics and Properties in Stress Corrosion Susceptibility",

NBSIR 75-916, October 1975.

3. Ibid., NBSIR 76-1170, November 1976.

4. M. Sakashita and N. Sato, Corrosion Science, 1_7, No. 6, 473 (1977).

5. K. Sugimoto and Y. Sanada, Corrosion Science, 17, No. 6, 473 (1977)

- 59 -
Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

XPS^peqtrogram o,f air-formed film on pure molybdenum.^ Plotted

are counts (2-1 OK) versus energy and shown are the Mo 3d peaks.

mm ovrszsq to sfoH sriT" <ed~8f PMM ,fbT6929.8 TovaM to sorttO

PlotSf pli-pea^ihpighfeirHfio data from Table 1 .

: - Shownsar-e Mb)/F^wcn3

for 4 surface conditions of both Fe-lMo and Fe-5Moj a-l loyls^ai^ 81281/1 wel 1 as Mo, v/M/ \ (m) (ox)

and Fe

(m)

/Fe

(0^

ratios.

'isdmovoil <0Uf-3\ 8I23H 4 .b’rdl XPS spectra for surface films following anodic passivation at

442 ; ftiV S{jE^wi$hou^ chloridfeiion;;' (solidoline/)fi2 ‘nhis ilpett^Qm ii *M

compared to the spectrum obtained for the air-formed film (dashed

)ine) onvFe^Mo. Kinetjc' energy inereasesbalong absc.issaiomrgu2

.f

.S

.£

.5

XPS spectrogram for surface film following anodic passivation

at 442 mV SHE with chloride ion (short dash line); this spectrum

is overlaid on the spectrum obtained for the air-formed film on Fe-5Mo (long dash line) for comparison.

Comparison of XPS spectra for Fe-lMo (solid line) and Fe-5Mo

(dashed line) in chloride solutions after anodic passivation

at 442 mV SHE.

-60 -

- Cd-
Figure 1

-61 -
Mo/Fe

Fe(m)/Fe(ox)

Figure 2

-6 2 -
INTENSITY

Figure 3

-6 3 -
Figure 4

-64 -
Figure 5

-6 5 -

NBS-1I4A (REV. t 1-771

U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBSIR 78-1429

2. Gov’t Accession

No.

3. Recipient’s Accession No.

4. TITLE AND SUBTITLE

The Role of Passive Film Growth Kinetics and Properties

in Stress Corrosion and Crevice Corrosion Susceptibility

5. Publication Date MAY 1978

6. Performing Organization Code

7. AUTHOR(S)

J. Kruger, J. R. Ambrose, J. J. Carroll, and A. J- Melmed

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 3120448

11. Contract/Grant No. NAONR 18-69, NR036-082

12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP) 13. Type of Report & Period Covered Office of Naval Research

Department of the Navy

Arlington, VA 22217

interim 12/76-12/77

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here.)

This report consists of the following four parts:

j _ Field Ion Microscopy Studies of the Interaction of Hydrogen with Selected

Metals and Alloys

II _ The Role of Noncrystalline Films in Passivation and Breakdown of Passivation

HI _ Effect of Chromium Depletion on the Repassivation Kinetics 18 10

Austenitic Stainless Steel in Sodium Chloride Solutions

IV - Composition of Surface Films Formed During the Repassivation of Iron and

Iron-Molybdenum Alloys

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a Rroper name; separated by semicolons) Crevice corrosion; field ion microscopy; iron; noncrystaili e

. . • i • j. • r-' 1 • nndniiim' Y-T*flV name; separatea uy aemit* j 1 c v ^ - - »- ^ films; passive films; repassivation; sensitization; stainless steel; uranium; x-ray

photoelectron spectroscopy

18. AVAILABILITY P0 Unlimited

For Official Distribution. Do Not Release to NTIS

1

1 Order From Sup. of Doc., U.S. Government Printing Offic Washington, D.C. 20402, SD Stock No. SNOO^-OOj

|X I Order From National Technical Information Service (NTIS)

Springfield, Virginia 22151

19. SECURITY CLASS

(THIS REPORT)

UNCL ASSIFIED

20. SECURITY CLASS

(THIS PAGE)

UNCLASSIFIED

21. NO. OF PAGES.

67

22. Price

$5.25

USCOMM-DC 6603S-P7E
.

'
i| H - 9 i. CD

I w

CD

) Oi

r ^

I cr

> 3

I P

pp

)

)

o H

9 3

C

ui §3

05 cu 3 £ 9 P

<i L

'J ^

OJo

CO

<1

00 S3

C/5

CDo.

o

?r

CD&

CD

r^- CDnrfo’-1

C/5

1 e I' 05 II 78-1430

Instrument : to Evaluate

ailed Smoke Detectors

k. Lee

ire Research

Applied Technology eau of Standards

. D C. 20234

rcui Udty I »78

Final Report

U S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS
..