Corrosion failures In service, there are many types of corrosive attack

12.2.2.3 Corrosion failures In service, there are many types of corrosive attack

which lead to rapid failure of components. A familiar example is intergranular corrosion and is associated with the tendency for grain boundaries to undergo localized anodic attack. Some materials are, however, particularly sensitive. The common example of this sensitization occurs in 18Cr–8Ni stainless steel, which

is normally protected by a passivating Cr 2 O 3 film after

heating to 500–800 °

C and slowly cooling. During cooling, chromium near the grain boundaries precip- itates as chromium carbide. As a consequence, these regions are depleted in Cr to levels below 12% and are no longer protected by the passive oxide film. They become anodic relative to the interior of the grain and, being narrow, are strongly attacked by the corrosion current generated by the cathode reactions elsewhere. Sensitization may be avoided by rapid cooling, but in large structures that is not possible, particularly after welding, when the phenomenon (called weld decay) is common. The effect is then overcome by stabilizing the stainless steel by the addition of a small amount (0.5%) of a strong carbide-former such as Nb or Ti which associates with the carbon in preference to the Cr. Other forms of corrosion failure require the compo- nent to be stressed, either directly or by residual stress. Common examples include stress-corrosion cracking (SCC) and corrosion-fatigue. Hydrogen embrittlement is sometimes included in this category but this type of failure has somewhat different characteristics and has been considered previously. These failures have certain features in common. SCC occurs in chemically active environments; susceptible alloys develop deep fissures along active slip planes, particularly alloys with low stacking-fault energy with wide dislocations and pla- nar stacking faults, or along grain boundaries. For such

selective chemical action the free energy of reaction can provide almost all the surface energy for fracture, which may then spread under extremely low stresses.

Stress corrosion cracking was first observed in ˛ -brass cartridge cases stored in ammoniacal envi- ronments. The phenomenon, called season-cracking since it occurred more frequently during the mon- soon season in the tropics, was prevented by giving the cold-worked brass cases a mild annealing treat- ment to relieve the residual stresses of cold forming. The phenomenon has since extended to many alloys in different environments (e.g. Al –Cu, Al –Mg, Ti –Al), magnesium alloys, stainless steels in the presence of chloride ions, mild steels with hydroxyl ions (caustic embrittlement) and copper alloys with ammonia ions.

Stress corrosion cracking can be either transgran- ular or intergranular. There appears to be no unique mechanism of transgranular stress corrosion cracking, since no single factor is common to all susceptible alloys. In general, however, all susceptible alloys are unstable in the environment concerned but are largely protected by a surface film that is locally destroyed in some way. The variations on the basic mechanism arise from the different ways in which local activity is generated. Breakdown in passivity may occur as a result of the emergence of dislocation pile-ups, stack- ing faults, micro-cracks, precipitates (such as hydrides in Ti alloys) at the surface of the specimen, so that highly localized anodic attack then takes place. The gradual opening of the resultant crack occurs by plas- tic yielding at the tip and as the liquid is sucked in also prevents any tendency to polarize.

Many alloys exhibit coarse slip and have similar dis- location substructures (e.g. co-planar arrays of disloca- tions or wide planar stacking faults) but are not equally susceptible to stress-corrosion. The observation has been attributed to the time necessary to repassivate an active area. Additions of Cr and Si to susceptible austenitic steels, for example, do not significantly alter the dislocation distribution but are found to decrease the susceptibility to cracking, probably by lowering the repassivation time.

The susceptibility to transgranular stress corrosion of austenitic steels, ˛-brasses, titanium alloys, etc. which exhibit co-planar arrays of dislocations and stacking faults may be reduced by raising the stacking- fault energy by altering the alloy composition. Cross- slip is then made easier and deformation gives rise to fine slip, so that the narrower, fresh surfaces created have a less severe effect. The addition of elements to promote passivation or, more importantly, the speed of repassivation should also prove beneficial.

Intergranular cracking appears to be associated with

a narrow soft zone near the grain boundaries. In ˛- brass this zone may be produced by local dezincifica- tion. In high-strength Al-alloys there is no doubt that it is associated with the grain boundary precipitate-free zones (i.e. PFZs). In such areas the strain-rate may be so rapid, because the strain is localized, that repassiva- tion cannot occur. Cracking then proceeds even though

Corrosion and surface engineering 387 the slip steps developed are narrow, the crack dis-

solving anodically as discussed for sensitized stainless steel. In practice there are many examples of intergran- ular cracking, including cases (1) that depend strongly on stress (e.g. Al-alloys), (2) where stress has a com- paratively minor role (e.g. steel cracking in nitrate solutions) and (3) which occur in the absence of stress (e.g. sensitized 18Cr–8Ni steels); the last case is the extreme example of failure to repassivate for purely electrochemical reasons. In some materials the crack propagates, as in ductile failure, by internal necking between inclusions which occurs by a combination of stress and dissolution processes. The stress sensitivity depends on the particle distribution and is quite high for fine-scale and low for coarse-scale distributions. The change in precipitate distribution in grain bound- aries produced, for example, by duplex ageing can thus change the stress-dependence of intergranular failure.

In conditions where the environment plays a role, the crack growth rate varies with stress intensity K in the manner shown in Figure 12.10. In region I the crack velocity shows a marked dependence on stress, in region II the velocity is independent of the stress intensity and in region III the rate becomes very fast

as K IC is approached. K ISC is extensively quoted as the threshold stress intensity below which the crack

growth rate is negligible (e.g. 10 ms ) but, like

the endurance limit in fatigue, does not exist for all materials. In region I the rate of crack growth is controlled by the rate at which the metal dissolves and the time for which the metal surface is exposed. While anodic dissolution takes place on the exposed metal at the crack tip, cathodic reactions occur at the oxide film on the crack sides leading to the evolution of hydrogen which diffuses to the region of triaxial tensile stress and hydrogen-induced cracking. At higher stress intensities (region II) the strain-rate is higher, and then other processes become rate-controlling, such as

Figure 12.10 Variation of crack growth rate with stress intensity during corrosion .

diffusion of new reactants into the crack tip region. In hydrogen embrittlement this is probably the rate of hydrogen diffusion.

The influence of a corrosive environment, even mildly oxidizing, in reducing the fatigue life has been briefly mentioned in Chapter 7. The S –N curve shows no tendency to level out, but falls to low S-values. The damage ratio (i.e. corrosion fatigue strength divided by the normal fatigue strength) in salt water environments is only about 0.5 for stainless steels and 0.2 for mild steel. The formation of intrustions and extrusions gives rise to fresh surface steps which form very active anodic sites in aqueous environments, analogous to the situation at the tip of a stress corrosion crack. This form of fatigue is influenced by those factors affecting normal fatigue but, in addition, involves electro-chemical factors. It is normally reduced by plating, cladding and painting but difficulties may arise in localizing the attack to a small number of sites, since the surface is continually being deformed. Anodic inhibitors may also reduce the corrosion fatigue but their use is more limited than in the absence of fatigue because of the probability of incomplete inhibition leading to increased corrosion.

Fretting corrosion, caused by two surfaces rubbing together, is associated with fatigue failure. The oxi- dation and corrosion product is continually removed, so that the problem must be tackled by improving the mechanical linkage of moving parts and by the effec- tive use of lubricants.

With corrosion fatigue, the fracture mechanics threshold K th is reduced and the rate of crack propagation is usually increased by a factor of two or so. Much larger increases in crack growth rate are produced, however, in low-frequency cycling when stress-corrosion fatigue effects become important.