Marcus P. Corrosion mechanisms in theory and practice
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universal models continue to appear from time to time and must be tested seriously.
Galvele’s surface mobility model [20,166] proposes that cracks grow by surface
diffusion of metal atoms (possibly combined with ions or molecules from the
environment) from a very sharp crack tip to the crack walls, i.e., capture of vacancies
by the crack tip (Fig. 42). Formation of surface compounds with low melting
points promotes SCC by increasing the surface self-diffusivity of the metal,
D
s
. The details of the crack tip and its interaction with the environment (i.e.,
electrochemistry) are considered to be secondary factors that contribute by
maintaining the critical surface compound. The model is applied to liquid metal
embrittlement as well as SCC and seems to have considerable predictive power.
One of the key factors in Galvele’s conversion to surface mobility was his
conviction that slip-dissolution could not work in the restricted geometry of the
crack, but, as explained earlier, the crack is not a one-dimensional slot and the
crack opening angle permits very high anodic current densities in cracks without
saturation of metal salts or very high IR potential drops.
442 Newman
Figure 41 Magnin’s model of SCC [16,17].
Copyright © 2002 Marcel Dekker, Inc.
According to Galvele’s model, the crack velocity (
υ
) is given by
Stress-Corrosion Cracking Mechanisms 443
Figure 42 Galvele’s surface mobility model for SCC [20].
where D
s
is the surface self-diffusion coefficient, Lis a characteristic diffusion length,
σ is the elastic surface stress at the crack tip, and a
3
is the volume of a vacancy. The
treatment of the diffusion process is greatly simplified, but this is not one of the main
flaws in the analysis. These are, according to Sieradzki and Friedersdorf [23]:
1. The equilibrium vacancy concentration in a plane can be increased by a
normal tensile stress [172], but Galvele applied the relevant equation to the
crack tip where this stress is zero. Lacking this effect, one must appeal to
small, second-order effects.
2. Proper accounting of the chemical potential of the vacancies must include
a capillary term of the form +γa
3
/r (γ = surface energy, r = crack-tip radius)
[173]. This prevents the crack from attaining very low r values, yet the
surface mobility mechanism assumes an atomically sharp crack.
Sieradzki and Friedersdorf end their analysis by concluding that the crack
velocities attainable by surface mobility (with the blunt crack tip just mentioned)
are 10
8
to 10
14
times lower than those calculated by Galvele. Their expression for
the crack velocity is
where N
s
is the number of lattice sites per unit area, Ω is the atomic volume (a
3
),
and κ is the curvature, which is just 1/r for the crack tip.
Copyright © 2002 Marcel Dekker, Inc.
SUMMARY AND RECOMMENDATIONS FOR THE FUTURE
Several models of SCC have been tested successfully in particular systems. It
would be premature to regard these tests as definitive (there is a particular need for
direct studies of crack dynamics in Au-Cu or similar alloys showing film-induced
cleavage), but they encourage us to look to other, unsolved systems for a more
complete picture of the range of SCC phenomena. Some of these systems are:
1. Creep-like phenomena that may involve cavity formation induced by internal
oxidation [60,61]. These will normally occur at elevated temperatures, and the
outstanding example is the SCC of alloy 600 in PWR primary water or
steam-hydrogen mixtures. The species being internally oxidized may include
both Cr and C [174]; hence Ni-C alloys may also crack under some
conditions. Some progress has been made by preexposing thin foils to the
SCC environment, then allowing hydrogen to escape and straining the foils
to failure at room temperature [61]; significant irreversible embrittlement has
been observed and is confined to the surface, suggesting an internal oxidation
effect rather than an irreversible hydrogen effect such as methane bubble
formation.
2. Gaseous embrittlement by species other than H
2
[101–103,141,142], where a
thorough test of the applicable models, using filmed thin foils would be
valuable.
3. SCC of iron or carbon steel under controlled potential conditions in
conducting, high-temperature aqueous environments [29], where more than
one mechanism may be operating, but the slip-dissolution and film-induced
cleavage models are eminently testable using established methods with
modern refinements.
4. Brittle SCC in iron exposed to CO-CO
2
-H
2
O [162,163] or anhydrous NH
3
[69]; indirect evidence for film-induced cleavage has been presented, but this
is too indirect at present, and the special surface layers that form in these
environments need to be studied.
5. Liquid metal embrittlement (LME) [15] (properly called metal-induced
fracture, as the metal need not be liquid so long as it is at a high T/T
m
)
remains an outstanding problem for SCC specialists. Lynch [19] has stated
that transgranular LME, in particular, involves only a surface interaction
with the environment and hence must be classed as an adsorption-induced
process even in the fcc materials; in situ electron microscopy, to observe the
cracking process directly, seems to be one way forward (although not with
such a crude environment as liquid mercury). The softening or local plastic
fracture claimed by Lynch may be demonstrated or may resolve into a
hardening or cleavage process at high enough resolution. Not enough is
known about the total crack tip system in liquid metal embrittlement, and a
proper accounting of all the forces and chemical potential gradients has not
been made.
6. Systems of interacting cracks should be studied using techniques developed
by statistical physicists [168].
Finally, there is room for more careful surface chemistry studies in SCC,
utilizing recent advances in the characterization of the crack solution (pH, etc.) and the
444 Newman
Copyright © 2002 Marcel Dekker, Inc.
crack path (segregation, etc.). Although these studies may not reveal anything radically
new, they are essential for a rigorous development of the subject, for example:
1. Interactions of anhydrous liquid ammonia with iron surfaces, to test the
proposed role of nitride layers [69,100]. An experimental problem is that
extremely low H
2
O levels are required to produce the pure Fe-N interaction
on a free surface, whereas in a crack the gettering action of the crack walls
maintains a very low water content at the crack tip.
2. Surface analysis of films formed transiently on bare, e.g., laser-irradiated
surfaces [175], should be undertaken. Such studies should include rapidly
solidified alloys that incorporate segregant elements and should be integrated
with mechanical studies on the same films.
3. The status of dealloying and film-induced cleavage in chloride-SCC of
austenitic stainless steels is still unclear. Further studies should be undertaken
in solutions that simulate the acid crack environment. The role of alloyed
P [88] should be examined using these or simpler (Fe-Ni) alloys.
4. High-resolution Auger or secondary ion mass spectrometry (SIMS) analyses
of SCC fracture surfaces, especially around crack-arrest features, have not
kept pace with instrument development, and there is great scope for
well-designed experiments of this type supported by cross-sectional
transmission electron microscopy.
5. The scanning tunneling and atomic force microscopes (STM and AFM) have
considerable potential for mapping fracture surface topography at very high
resolution [176], provided one chooses an appropriate system with minimal
corrosion and large flat areas on the fracture surface—this is illustrated every
time one tests an STM or AFM on cleaved graphite or mica. Such studies
may resolve some of the micromechanics of film-induced cleavage—for
example, by showing that crack-arrest markings are present everywhere on
transgranular fracture surfaces but are simply too fine to be resolved in the
SEM. One may also be able to use the new generation of these instruments,
with their improved position control and microscope facilities, to carry out
high-resolution side-surface studies during SCC.
Finally, while SCC of metals is important in our industrial society, SCC of other
inorganic solids can be said to have shaped our whole world. It is now widely accepted
that time-dependent fracture (stress corrosion) of rocks in contact with water is
responsible for such phenomena as premature fault movement leading to earthquakes
[177] and fracture of newly formed crustal rocks emerging on the ocean floor [178].
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London, 1976, ch. 8.1.
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175. R. Oltra and M. Keddam, in Ref. 68, p. 17.
176. J. Lankford and M. Longmire, J. Mater. Sci. 26:1131 (1991).
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178. P. G. Meredith (University College, London), private communication.
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Copyright © 2002 Marcel Dekker, Inc.
13
Corrosion Fatigue Mechanisms in Metallic
Materials
T. Magnin
Ecole Nationale Supérieure des Mines de Saint-Etienne, Saint Etienne, France
INTRODUCTION
The deleterious effect of an aqueous environment on fatigue crack initiation and prop-
agation in metals and alloys has been observed for a long time [1–5]. The applied elec-
trochemical potential has a large influence on localized corrosion reactions and, con-
sequently, on corrosion fatigue (CF) damage. These complex effects are used for
cathodic protection in offshore structures. Nevertheless, CF is influenced by various
mechanical, chemical, and microstructural parameters that interact locally. Even if
anodic dissolution versus hydrogen effects are known to occur during CF, the damage
mechanisms are still controversial and fatigue lifetime predictions are still empirical.
The application of linear elastic fracture mechanics (LEFM) led to the
determination of threshold stress intensity factors for long crack growth, which are
quite useful for engineering material–environment systems. Nevertheless, it is now
well established [6,7] that short cracks can develop and grow from smooth surfaces,
due to localized corrosion-deformation interactions, at crack tip stress intensities
below the previous threshold. CF damage models of short crack growth are then
necessary to improve the fatigue damage predictions in aqueous solutions.
In the first part of this chapter, crack initiation mechanisms are reviewed and
quantitative approaches are discussed, bearing in mind the limitations related to the
evaluation of localized corrosion-deformation interactions. A particularly interesting
example of electrochemical and mechanical coupling effects during CF in duplex stain-
less steels will illustrate the complexity of the interactions that lead to crack initiation.
Crack propagation models are then presented for both short and long cracks.
Anodic dissolution and hydrogen effects at the crack tip are analyzed. Finally, the
possible relation between stress corrosion cracking and CF is shown for crack
initiation and crack propagation processes.
CORROSION FATIGUE CRACK INITIATION
It is well known that slip bands, twins, interphases, grain boundaries, and constituent
particles are classical sites for crack initiation. Moreover, persistent slip band
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Copyright © 2002 Marcel Dekker, Inc.