Stress Corrosion Cracking

3 Stress Corrosion Cracking J. C. SCULLY

I. Introduction There is a long history of alloy failures that have been induced by the presence of a corrosive environment. One of the most serious forms of failure, observed over a period of time stretching back to before the beginning of this century, has been stress corrosion cracking, a phenomenon by which a highly localized form of corrosion causes some alloys to develop cracks when subjected to stresses that may be well below the value of the 0 1 % Proof Stress. In service failures the source of the acting stress is frequently residual, generated by fabrication procedures, e.g. welding, or forming processes, e.g. drawing. In laboratory investigations the same source may be used but more commonly the stress is usually applied externally since it is then easy to control and reproduce. Under operating conditions a stress on a component in some structures may not be constant. A vessel or a pipeline, for example, may be periodically filled and emptied. The subject must therefore also include reference to effects caused by a fluctuating stress since this may, on the one hand, exacerbate stress corrosion effects and, on the other, result in a different type of fracture, i.e. corrosion fatigue. Stress corrosion cracking occurs in many alloys and it has been the cause of a large number of service failures, many minor, some major, particularly in the chemical and transport industries. It is also of considerable importance in the nuclear power industry in which, for example, conditions may exist in which austenitic stainless steels may fail intergranularly in high purity pressurized water containing

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dissolved oxygen and chloride ions at the level of pbb. Because of its widespread importance stress corrosion cracking has been widely investigated and reported in a range of conference proceedings (for example, Dix, 1945; Robertson, 1956; Rhodin, 1960; Staehle et ai, 1969; Scully, 1971; Brown, 1972; Swann et ai, 1977; Staehle et ai, 1978; Foroulis, 1979; Arup and Parkins, 1979), books (for example, Logan, 1966) and reviews (for example, Parkins, 1964; Parkins, 1972; Scully, 1973; and many in the cited conference proceedings), in addition to an immense number of papers in the open scientific literature and an unspecifiable number of internal organizational reports. The history of the subject goes back to the last century in which reports were concerned with the season cracking of cold-drawn brass cartridges in ammoniacal atmospheres and early in this century with the caustic cracking of riveted boilers. Since then an increasingly wider perception of the problem has been forced upon the engineer. Austenitic stainless steels and other microstructural forms of stainless and non-stainless steels, aluminium, magnesium, titanium and zirconium alloys have also provided examples of susceptibility and, at the present moment, it is not possible to specify any system of commercial alloys in which at least some alloys are either not susceptible or unlikely to be so although not, perhaps, yet reported. Before, however, too pessimistic a picture is painted it is important to emphasize at this early stage that several distinctions need to be made. Failures observed in laboratory studies may not always be reproduced under service conditions. The degree of susceptibility in practice also becomes an important consideration: How easily does an alloy crack? This point is considered further below. Only some alloys in any system are susceptible and in some cases then only after a particular heat treatment, for example. From an early stage it was recognized that stress corrosion cracking occurred only in certain environments. Table 1 gives a list of the principal combinations of alloys and environments. Increasingly it has become apparent that the specific nature of the environment for any alloy has become less apparent than was earlier thought. Admiralty brass, for example, can fail in non-ammoniacal environments under open circuit conditions, e.g. tartrate solutions (Johnson and Leja, 1966). Table 1 is not intended to be comprehensive with respect to environment or to alloys. Each few years, it seems, some new phenomenon of stress corrosion cracking is observed: an alloy previously considered immune or an environment thought to be harmless must be added to the list. Clearly there must be limits but these have not yet been

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reached. As the requirements made upon alloy strengths increase and the environments to which they are exposed increase in composition, pressure and temperature, the incidence of newly encountered, unanticipated failures can be expected to continue. In this chapter the subject of stress corrosion cracking is discussed in several sections, with the emphasis being on the underlying mechanisms. The methods of testing are first briefly discussed since it is important to the subsequent description to appreciate that different methods may reveal different results. Secondly, the problems associated with the breakdown of protective films on metal surfaces are described, with particular emphasis upon the repassivation process and the significance of its kinetics. Having established the conditions under which the metal may be unfilmed, the possible reactions between the metal and solutions are then discussed. Some alloy systems are then examined: austenitic stainless steels, 70Cu-30Zn alloys, titanium, aluminium and zirconium alloys. Finally the fracture morphology is discussed together with a detailed consideration of events occurring at the crack tip, including the influence of factors already described in the previous section. It is intended that a balanced modern picture of the stress corrosion process should be provided particularly with a view to drawing systems together in order to provide a general picture where this is helpful. Inevitably, with such a large subject and with a continuing output of new work, changes are occurring continually. Stress corrosion cracking occurs as the result of an interaction of metallurgical, mechanical and electrochemical or chemical kinetic TABLE 1 Alloy/environment systems exhibiting stress corrosion cracking Alloy

Environment

Mild steel

Hot nitrate, hydroxide and carbonate/bicarbonate solutions Aqueous electrolytes, particularly when containing H 2 S Hot, concentrated chloride solutions; chloridecontaminated steam High purity steam Ammoniacal solutions Aqueous, Cl~, Br~ and I" solutions Aqueous Cl~, Br~ and I" solutions; organic liquids; N 2 0 4 Aqueous Cl" solutions Aqueous Cl" solutions; organic liquids; I 2 at 350°C

High strength steels Austenitic stainless steels High Ni alloys α-Brass Al alloys Ti alloys Mg alloys Zr alloys

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events. Any appreciation of mechanistic points must include each of these events. Only under very specific conditions do these various rate processes combine to give stress corrosion crack propagation. Only a qualitative description is possible in most cases but there does appear to be enough of a common description to most processes of stress corrosion cracking to expect that, in the next few years, detailed quantitative models will be developed. II. Testing for stress corrosion cracking A. Mechanical

arrangements

A variety of experimental tests has been employed in order to investigate stress corrosion cracking, many of which are described in an ASTM book (Champion, 1967) and by Parkins et al. (1972). Before describing these various testing methods it is necessary to point out both that the choice of a test and the development of variations of a specific type of test will be determined by the principal objectives of the test. A successive series of questions and answers arise: Does the alloy under consideration crack in the environment in which it will be exposed in service? This would appear to be the first and simplest question. The possibility of failure is the first detail that must be established. In order to determine this a clear positive test is required. What are the chances of failure? More difficult, and equally necessary, is the requirement for a second type of test to follow upon the first if possible failure has been established. This should give an indication of probability, since in order to avoid stress corrosion failure it may only be necessary to lower the probability rather than achieve the more complicated task of achieving complete immunity, which in most cases will involve changing the alloy, employing inhibitors etc. In many examples such changes may involve considerable additional expenses if designs have to be changed or operating procedures modified. If the probability of failure is low, however, such expensive remedial actions may be unnecessary. Many alloys are used in environments which, under somewhat different conditions, can cause stress corrosion cracking. In the first instance, the possibility of failure is the first consideration. This may resolve itself into testing the actual component in the environment to be used under service conditions. Once an attempt is made to accelerate the test, by raising the applied stress to values higher than those encountered in practice, or by intensifying the

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environment beyond the predetermined range of operating conditions by, for example, raising the temperature or concentration, problems in interpretation are likely to arise. If rapid failure occurs under such accelerated conditions it can be difficult to extrapolate such results to the less intense conditions that will be encountered in service, and that problem is recognized. The reverse problem can also arise with regard to environmental conditions. Immunity over a certain temperature range does not automatically eliminate cracking occurring at a lower temperature range even if in most cases this is observed. Tests can be grouped under several headings, depending upon the mechanical variable employed: load, constant strain, the use of stress intensity and constant crosshead speed. Constant load tests are widely employed, often on tensile specimens, and give information of the type shown schematically in Fig. 1 in which, following the common practice, the initial stress value is plotted rather than the load. Two simple points can be made about the results. First it may be difficult to ascertain whether a threshold stress exists. Scatter in such tests can be large, particularly at high values of the time-to-failure, tf. Secondly, on smooth uncracked specimens, the measured value of t{ consists of two components, the time-to-initiate, t{, and the time-to-propagate, /p, where: h + tp = tf.

(1)

ω I

o

"7 0)

E

Initial stress

FIG. 1. Schematic diagram of the common relationship observed between the initial stress and time-to-failure in constant load tests.

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If this type of test is used to examine any metallurgical or environmental variable, then it may not be clear which component is being affected the more by changes in a variable, or whether one is totally unaffected. This point may not be of great importance, particularly where t, » /p, and this type of test is widely employed in many laboratories in examining both metallurgical and environmental variables. Constant strain tests employ specimens that are bent in a reproducible manner and clamped into shape, e.g. U-bend specimens. Time-to-failure may be measured although often such specimens exhibit a degree of stress relieving when crack propagation occurs and total failure may not then occur. The time to observe the first development of visible cracks may then be the measured effect. Such tests are very cheap when compared to other types of test and they are commonly employed as a general screening test. The use of tests employing stress intensity measurements with precracked specimens arises from the general recognition that engineering structures always contain surface cracks and flaws which may exhibit a wide range of depth and acuity. In relation to equation (1) the question then arises under what conditions such a crack or flaw propagates, because initiation conditions are rapidly achieved in such tests and in many situations ti « tp. Over the last 20 years fracture mechanics has been applied increasingly to stress corrosion cracking (Brown, 1968). In the simplest type of test the time-to-failure is measured as a function of initial stress intensity factor, as shown in Fig. 2. Where a threshold value is obtained this is designated Klxc where I refers to the opening mode (Knott, 1973). The ratio KlsJKlc is a useful comparison standard of susceptibility. Where the ratio is >0·9 susceptibility is low whereas where the ratio is <0·5 susceptibility is high. Titanium alloys in sea water exhibit a threshold stress intensity (Brown, 1972) but this is not seen in many systems. Various experimental arrangements have been employed, as described by Brown (1972). An extension of the presentation shown in Fig. 2 has been to measure the crack propagation rate as a function of the stress intensity factor for stress corrosion systems, as had been done previously by Wiederhorn for glass in water (1968). For high strength steels and aluminium and titanium alloys a general pattern has been observed as drawn schematically in Fig. 3. At low values of K, greater than Khcc where this is observed, the crack propagation rate increases logarithmically with increasing K. At some point as K increases the propagation rate reaches a maximum or plateau value at which the rate becomes mass transfer controlled. It is then dependent

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Time-to-failure

FIG. 2. Schematic diagram of time-to-failure as a function of initial stress intensity factor for two systems, one of which exhibits a threshold stress intensity, Khcc, below which failure does not occur.

upon the diffusion of reactants or products to or from the crack tip. The rate, for example, is lowered by an increase in solution viscosity. Stage III is observed in some alloys but usually it is not observed so that stages I and II often constitute the complete graph. Such data can be of value in alloy selection as well as in the elucidation of propagation mechanisms. Constant crosshead speed tests have been employed by a number of workers over the last 20 years or more but in the last 5 or 10 years its employment has increased considerably (Ugiansky and Payer, 1979). Specimens are subjected to low tensile strain-rates (10~ 7 -10 - 4 s - 1 ) while in contact with an environment. Susceptibility is assessed by comparison of values of reduction in area or elongation-to-failure obtained in the environment to those obtained in an inert environment at the same temperature as the stress corrosion test. Measurements and calculations of crack velocity are also sometimes employed (Scully, 1979). This type of test is relatively rapid and quite severe. By imposing low strain-rates on metal surfaces it reproduces a mechanical condition similar to that existing at the crack tip surface during crack propagation.

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Stress intensity factor,

K

FIG. 3. Schematic diagram of stress corrosion crack propagation rate in high strength alloys as a function of stress intensity factor. In many alloys stage III is not seen.

B. Electrochemical

arrangements

On the environmental side of testing, conditions must be imposed that are of immediate relevance to the problem under investigation. Much work is done under open circuit conditions but in looking for susceptibility it is important to realize that cracking may occur over a range of potential which does not include the corrosion potential. An example is shown in Fig. 4 (Uhlig and Cook, 1969) which shows results, obtained for an austenitic stainless steel exposed to MgCl2 solutions boiling at 130°C, of time-to-failure as a function of potential. While the addition of 2 wt. % N a N 0 3 to the solution narrows the range of potential over which cracking occurs, the corrosion potential falls within the narrowed range and cracking occurs under open circuit conditions. With the addition of 5 wt. % N a N 0 3 to the solution the range is further narrowed while the

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STRESS CORROSION CRACKING

corrosion potential now lies outside it. Cracking will not therefore occur under open circuit conditions. It does occur, however, if a small amount of anodic polarization is applied. Unless the role of potential is examined it might be quite erroneously concluded that 5 wt. % of the inhibitor removed susceptibility completely. The simple conclusion that the absence of failure under open circuit conditions does not imply the absence of susceptibility merely illustrates the importance of examining the role of potential in controlling cracking, analogous, perhaps, to the obvious importance of examining a range of values of stress. Electrochemical methods of predicting susceptibility to stress corrosion cracking have been examined in detail at a 1978 Conference (Scully, 1980a, b). One particular method that has shown some success has been a potentiodynamic scan-rate method. An anodic polarization curve obtained at a rapid scan-rate is compared with the same curve obtained at a low scan-rate. Where film forming

30 h NaNO-, - j_ *- COrr

>

1 2 % NaN0 3

E

- I

x

CO

kj

90

150

^corr

MgCU + 2 % NaN0 3

I

5% NaN0 3

10

100

Time-to-failure ( h )

FIG. 4. The relationship between applied potential and time-to-failure for cold rolled 18Cr-8Ni steel in MgCl2 boiling solutions at 130°C to some of which additions of 2 and 5 wt. % N a N 0 3 have been made. The corrosion potentials are also shown (Uhlig and Cook, 1969). The addition of 5 wt. % resulted in an open circuit potential that was outside the range of potential in which cracking occurred. Without conducting tests at different potentials it might be concluded quite erroneously that susceptibility had been completely removed.

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conditions are relatively slow the rapid scan-rate will show a large dissolution rate at certain regions of potential where a slow scan-rate will show a low anodic current because there has been time for film formation to occur. By inspection of the two curves it becomes possible to discern regions of potential where cracking might occur and to estimate the likely maximum crack propagation rate. The technique has been applied to ferritic steels in a number of environments with success although there are limitations upon its wider use. The film formation rate, for example, must be relatively slow and if the solution pH changes appreciably during the polarization there may be additional problems in interpretation. An example of such results is shown in Fig. 5 which is taken from the extensive work of Parkins (1980) who has developed this approach to prediction. The predicted and observed potential ranges for cracking are in reasonably good agreement. While there is some element of subjectivity involved in such interpretations Parkins (1980) reports that cracking conditions are reasonably accurately defined if either of two boundary conditions are applied: (a) that current densities < 1 m A c m - 2 on the fast scan-rate are regarded as negligible, and (b) that the current density difference between fast and slow scan-rates, Δ/, divided by the current density indicated by the slow scan-rate, is, at a given potential, exceeds 1000. Anodic protection

-400

I Stress I corrosion

-600

Pitting

-800

Icathode protection

-1000

Potential (mV(SCE)) FIG. 5. Potentiodynamic polarization curves for a C-Mn steel in 1 N N a 2 C 0 3 + 1 N N a H C 0 3 solution at 90°C showing the domains of behaviour predicted from the curves (Parkins, 1980).

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III. General mechanistic features A. Film

breakdown

Stress corrosion is often considered to be the most highly localized form of corrosion. To achieve this the metal lattice and the solution must react in a particular way. Before considering this reaction and its many ancillary features, the nature of the metal surface must be taken into account. In most cases the metal and the environment are not initially in direct contact with each other because the metal is filmed. Most metals and alloys in aqueous solutions develop a surface film that is characteristic for the solution and is often different from the air-formed film existing prior to immersion. While it is obvious that passive alloys are filmed, e.g. stainless steels, it is also true for mild steel and copper alloys on which the film may have characteristics different from a passive layer, e.g. it may be comparatively thick or have a very high conductivity. Included in such a general description would be thin protective metal films that may develop from a dealloying phenomenon, e.g. copper or copper-rich films on a-brass. Before any form of corrosion reaction occurs between the bulk alloy and any of the constituents of the solution the two must come into contact. This can only happen if the surface film, whatever its nature, is removed locally either by chemical or by mechanical means. Chemically, film breakdown may occur by the penetration of film pores as is described in the chapter on pitting (Galvele, this volume, p. 29) if the corrosion potential is above the breakdown or pitting potential or, in the case of non-passive film-forming alloys, below or outside the potential range over which the particular film ceases to be stable. Mechanically, the breakdown process in stressed specimens is commonly thought to be caused directly by the action of the stress in producing a surface strain-rate, albeit very low, and causing plastic deformation. A very simplified diagram is drawn in Fig. 6 in which an emergent slip step is drawn as intersecting and breaking a surface film, thereby revealing fresh metal which will react with the solution (Scully, 1967). Such a drawing is intended to represent a very simple picture of what is likely to be an event that for a number of reasons is rather more complicated. For example, the mechanical properties of the film are clearly going to be important since, in the extreme case, if the film were very ductile it might deform and stretch and not cause the exposure of any fresh metal surface. In addition, the thickness of the film, t, in relation to the height of the slip step, A, will be important. If tjh is relatively high, fresh metal may not be revealed at

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all. Thus thick films may be more beneficial than thin films, provided that their mechanical properties in the two forms are unchanged. Alloys exhibiting lamellar slip because of relatively easy cross slip are likely to be less susceptible because the average slip step height will be low.

C

D

FIG. 6. Schematic diagram of a freshly created slip step breaking a filmed surface at the tip of a stress corrosion crack. A. Completely filmed surface. B. Slip step fractures film. C. Corrosion occurs on the freshly created metal surface as well as film formation. D. The surface is repassivated after the extension of the crack.

These factors of film ductility and slip step height have been discussed at length previously (Swann, 1963). While it has proved difficult to show unequivocally whether the relative ductility of films has any significance in stress corrosion cracking mechanisms there is evidence that a film affects the behaviour of the metal. The presence of a film on nickel, for example, does affect the slip mode of the subjacent metal (Latanision and Staehle, 1969). The slip mode does appear to be directly relevant, however, It is characteristic of FCC alloys (Swann, 1963), and perhaps of hexagonal alloys also (Sanderson and Scully, 1967), susceptible to transgranular cracking, that they exhibit co-planar arrays of dislocations, indicative of a tendency not to cross slip easily, when examined by transmission

STRESS CORROSION CRACKING

115

electron microscopy. While the original explanation for this observation in FCC alloys centred around the significance of a low stacking fault energy (Swann and Nutting, 1959/60), such co-planar arrays may arise because of the occurrence of short-range order and for other unspecified reasons (Swann, 1963). The tendency to form coplanar arrays of dislocations, which can be interpreted as giving rise to large slip steps and high stresses at grain boundaries at points where pile-ups impinge, appears to be a necessary requirement for the occurrence of transgranular cracking in single phase alloys. In such a complex phenomenon as stress corrosion cracking, it is perhaps not surprising that it is not a sufficient requirement. Many alloys exhibit such dislocation arrangements without being susceptible to stress corrosion cracking. The additional requirements for susceptibility are described below. When the situation shown in Fig. 6 develops, the subsequent reactions occurring on the freshly exposed metal step can be supposed to be of considerable complexity. Only a simple description is attempted, at least initially. First, the alloy will start to dissolve. Secondly, H 3 0 + discharge may occur, with a fraction of the H atoms formed being absorbed by the metal. Thirdly, the film will start to form on the new surface, since a filmed surface is the stable configuration of the alloy surface. The overall rate of each of these reactions, which are subdivisible into partial reactions, will depend upon the value of the electrochemical potential of the new surface and upon local hydrodynamic considerations, stemming both from the rapid movement of the surface and from gas evolution, both of which will contribute to an overall general stirring effect upon the volume of liquid at the crack tip. In addition, the occurrence of several reactions on sites immediately adjacent to each other may result in secondary reactions between reactants and products of one reaction with another. It is not possible to pursue this description further since data about such events are difficult to obtain but it is important to emphasize that the simple picture conveyed in Fig. 6 in reality is likely to be much more complicated. Developments in the solution at the crack tip must also be considered. An occluded cell develops. The formation of metal cations and their subsequent hydrolysis leads, via a number of intermediate steps, to: Me z+ + Z- H 2 0 ^± Me(OH)z + zH + .

(2)

Anion migration into the region of the crack tip from the bulk

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solution occurs in order to maintain charge neutrality. In chloride solutions, for example, the pH measured at the region of the crack tip corresponds to that of the saturated metal chlorides: pH 3-8 for steel, 3-5 for aluminium and 1-7 for titanium (Brown et ai, 1969). These values are similar to values observed inside pits in the same materials. The first two reactions of dissolution and hydrogen ion discharge are responsible jointly or singly for the crack propagation process. It has proved very difficult to obtain a clear picture for any particular alloy of which one is the principal reaction or to explain why dissolution should be so localized. Section B on Mechanisms describes these difficulties in detail but what must be emphasized at this stage of the description is that these two reactions can occur only as long as the freshly exposed metal is "bare" or at least incompletely filmed. Once film formation has advanced sufficiently, such reactions will be either arrested or reduced to an insignificant rate. As the supply of cations falls the maintenance of the occluded cell conditions will no longer be possible. Both the anion concentration and hydrogen ion concentration will start to fall. The important reaction of film formation has been called repassivation (Scully, 1967) and much emphasis has been laid upon it, as is described by Ambrose (this volume, p. 175). It has been argued, for example, that the repassivation rate is of critical importance in the stress corrosion crack propagation process (Scully, 1967). If the single slip step event of Fig. 6 is depicted electrochemically as shown in Fig. 7, the current on the emerging step will rise to a maximum from which it decays as repassivation occurs. The single transient represents a flow of charge Jf0 / dt. In practice such single events cannot be easily followed although the occurrence of a large number of such events is readily detectable (Hoar and Scully, 1964). In measuring such currents attention must be given to hydrogen evolution since if this occurs simultaneously on emergent steps it will reduce the measured anode current by an equivalent amount. Overall it is the amount of corrosion that is important. If repassivation is relatively rapid then insufficient charge will flow to sustain cracking and crack arrest will occur. If repassivation is relatively slow then corrosion will not be localized but will spread laterally, causing crack blunting and resulting not in a crack but in an elongated fissure or pit. Such an argument determines that cracking occurs with a repassivation rate within some range, possibly quite narrow, of values (Scully, 1967). Where hydrogen absorption is responsible for the crack propagation process rapid repassivation is also likely to retard or reduce cracking since the formation of the film

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STRESS CORROSION CRACKING

will hinder hydrogen entry and the hydrogen discharge rate. If repassivation is very slow, however, this will not affect cracking since it cannot be supposed that too much hydrogen can be absorbed over too great a surface. The proposed range of repassivation values is therefore much less restricted in stress corrosion mechanisms incorporating hydrogen embrittlement processes. Many inhibitors reduce or eliminate stress corrosion by increasing repassivation rates. At the crack tip the ratio of inhibitor/chloride ion is likely to be of critical importance (Scully, 1968). Organic acid inhibitor molecules are often relatively ineffective in this respect, one reason being their low diffusion rates which make it difficult to maintain a sufficient concentration at the advancing crack tip. Film fracture

Film repair

Passive current

Time

FIG. 7. Transient current generated by the process depicted in Fig. 5 as a function of time. The area under the curve is the charge that flows.

If the range of repassivation rates is narrow when dissolution is important then it can be expected that the value of the potential at the crack tip over which cracking might occur will occupy a narrow range of values since the repassivation rate is very dependent upon potential (Engseth and Scully, 1975). If repassivation is too slow pitting occurs. The cracking potential range is just below the pitting potential in a number of alloys. The active-to-passive potential transition region is also a region where cracking might be expected since the rate will be changing rapidly in that region. These two ranges are drawn in Fig. 8, and examples of both can be observed in the literature (Staehle, 1971).

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The electrochemical reactions governing repassivation must be considered in conjunction with the mechanical and metallurgical reactions that create the fresh metal area continually as the crack propagates, a process that is caused by the crack tip strain-rate. What Fig. 6 does not convey is the dynamic aspect of the slip step creation process. Taken together the propagation process can be considered as arising from an imbalance between a repassivation process and a creep process. An early illustration of this is depicted in the results shown in Fig. 9 for a Ti-5Al-2*5Sn alloy exposed to two different types of solution in a constant extension rate test (Scully and Powell, 1970). In aqueous NaCl solution the alloy is passive. Stress corrosion cracking occurs over a narrow range of crosshead speeds. At the upper limit ductile failure occurs more rapidly than crack initiation.

Cracking zone

Cracking zone

Log current density

FIG. 8. The regions of potential where stress corrosion crack propagation may occur in relation to a schematic anodic polarization diagram for an alloy exhibiting an active-to-passive transition. In the two areas indicated the repassivation rate will undergo a considerable change in value since, in both cases, the surface is changing from passive to active.

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At the lower limit repassivation occurs more rapidly than fresh metal area is being created. In the CH 3 OH/HCl mixture the alloy corrodes. Stress corrosion cracking occurs over a much wider range of crosshead speeds. At the upper limit ductile failure occurs for the same reason as in the aqueous solution: there is no time for crack initiation. As the crosshead speed is lowered cracking continues to be observed below the value where in the aqueous solution the alloy would fail to crack because in the corrosive mixture repassivation is not possible. The lower boundary of crosshead speed will depend upon the repassivation characteristics of the system. This will be determined by the solution composition, specimen potential and alloy composition. Tests on cracking specimens have shown that crack arrest can occur as the crack tip strain-rate is lowered (Scully and Adepoju, 1977). This value has been designated έΓ (Scully, 1975). The important emphasis is on the repassivation process being able to give complete protection not only to a static surface area but also to a surface area that is being continually increased. The conditions under which this occurs are the result of a specific interaction between two 14,

1

-1-0

-0-5

0

0-5

Log crosshead speed (mm min-1)

FIG. 9. The relationship between elongation to fracture and crosshead speed for a Ti5Al-2-5Sn alloy in 3 % aqueous NaCl solution and a CH 3 OH/HCl mixture. In the aqueous solution repassivation occurs and prevents stress corrosion cracking as the employed crosshead speed is lowered. In the CH 3 OH/HCl mixture no such repassivation is possible and cracking continues to be observed as the employed crosshead speed is reduced. (Scully and Powell, 1970.)

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kinetic processes, a strain-rate and a repassivation rate, and much attention has been focused upon this relationship (Scully, 1980). The importance of film fracture and repair has long been emphasized. Between temporarily unfilmed metal at the crack tip and passivated crack tip sides a substantial potential difference (500 mV) may exist in some systems, e.g. Fe in boiling N H 4 N 0 3 solutions (Logan, 1952). Logan attached much significance to such differences. Related ideas have come from the work of Staehle (1971), Engell (1971), Parkins (1972), Vermilyea (1972) and Bignold (1972). The overriding control of film formation on crack propagation processes can scarcely be overemphasized. It is a particularly important factor to be borne in mind when attempts are being made to elucidate stress corrosion cracking mechanisms, by, for example, variations of environmental factors. Any such variation, in potential or pH, for example, will affect the metal/environment reactions that are discussed in the next section but, in addition, it will also affect repassivation rates. Mechanistic conclusions about particular effects have therefore to be drawn with great care. Changes in repassivation rate may have a greater and in some cases an apparently contrary effect than anticipated or inferred changes in the metal/environment reaction. While such a description is a simple one it may not be easy to discern what the cause of any particular observed effect is. Cathodic polarization, for example, of a passivatable alloy, will usually promote repassivation and thereby arrest cracking, if cracking is occurring in the more noble region of potential in Fig. 8. Such a result may reveal nothing about the mechanism of cracking since such arrest would be expected independently of how the crack propagates. Another simple example arises from the common observation that many alloys exhibit a range of potential over which they crack, an example of which has already been given in Fig. 4. If the corrosion potential falls outside the cracking range, as was the case for additions of 5wt. % N a N 0 3 , the alloy appears to be non-susceptible under open circuit conditions. If, for example, anodic polarization induces cracking it does not indicate that dissolution is the principal cracking mechanism. The function of such polarization is to produce film breakdown (and lower the repassivation rate). These two examples are relatively simple and perhaps seemingly obvious. The exact interpretation of the effect of altering environmental factors is, however, a very demanding task and the literature contains numerous examples of definitive claims which cannot be substantiated, e.g. that cathodic protection of titanium alloys in sea water illustrates that hydrogen embrittlement is not responsible for the stress corrosion

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crack propagation process (Feige and Murphy, 1966), whereas, instead, it illustrates the effect of a higher repassivation rate in a situation where the passive film is highly impermeable to hydrogen. B. Cracking

mechanisms

In the preceding section much emphasis was placed upon the role of the surface film in controlling the crack propagation process. During propagation the metal surface at the crack tip is either continuously or intermittently unfilmed while the crack sides are permanently filmed. Such a description would explain why the corrosion reaction is confined to the crack tip, occurring in a series of transients, one of which is drawn in Fig. 7. While one part of the reaction of the newly generated metal surface and the environment will be the film repair process itself, the other part will be the process that causes crack extension. This latter part has always received considerable attention since it lies at the root of the answer to the fundamental question: Why does crack extension occur? Two very general ideas have been put forward, both of which can be subdivided into several categories. The first is concerned with the anode reaction at the crack tip. As a result of dissolution the crack extends, a mechanism often described as Active Path Stress Corrosion Cracking. The second is concerned with the cathode reaction at the crack tip. As a result of local hydrogen absorption and consequent embrittlement the crack extends, a mechanism often described as Hydrogen Embrittlement Stress Corrosion Cracking. 1. Active path stress corrosion cracking This is the older of the two mechanisms and stems from regarding stress corrosion cracking as the extreme example of localized corrosion. Cracking extends as a result of anodic dissolution because of the existence of an active path, a narrow region of metal that will dissolve rapidly. Why dissolution should occur on such a fine scale is either because of a pre-existing path, e.g. a grain boundary containing impurities, as a result of which a cell is created causing local dissolution at a rate significantly higher than the matrix, or because the imposed stress causes a path to be generated, commonly described as a strain-induced path. The effect of a pre-existing path may lie not so much in the relative dissolution rates of the grain boundary and matrix as in the relative rates of repassivation. If the matrix repassivates and the grain boundary does not, then large differences in corrosion rate between

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the two can be expected to occur. The function of the stress must be considered also, since intergranular stress corrosion crack propagation does penetrate more rapidly than intergranular corrosion (e.g. Majumdar et al, 1980). The imposed stress may serve merely to break the film forming at the grain boundaries for the reason drawn in Fig. 7 so that the average current causing grain boundary separation is higher through being continually broken. If the film forms by precipitation of dissolved metal it may be thicker over the more active grain boundaries. Such effects are seen in ferritic steels in nitrate solutions (Parkins, 1971) and in α-brasses exposed to alkaline ammoniacal solutions (Pugh, 1971). The existence of a pre-existing path can be inferred from a number of experiments although the actual cause in each case is not always clear. Unstressed specimens of α-brasses, for example, undergo intergranular corrosion in 15 N ammoniacal solutions in which they become covered with a thick black porous film (Pugh, 1971). Very little is known about the composition of the grain boundaries of abrass and why they should corrode at a higher rate than the matrix cannot be explained. When a stress is applied cracking becomes transgranular in 70Cu-30Zn alloys (Kermani and Scully, 1979) which indicates that the pre-existing path may not be the most reactive path once a stress is applied. In 70Cu-29Zn-lSn alloy the cracking in the same solution is intergranular (Pugh, 1971), however. The effect of alloying difference in changing the predominant crack path cannot be explained. Application of prior cold work to the 70Cu-29Zn-lSn alloy gives rise to transgranular cracking (Pugh, 1971). In the absence of an explanation, all that can be supposed is that there are two competing crack paths, possibly associated with different causes, one of which is pre-existing while the other is generated by plastic deformation. Which occurs preferentially will depend upon local chemical reactions, surface chemistry and compositional effects and the influence of dislocations. A similar situation is observed in Type 304 Austenitic Stainless Steels exposed to 5 N H 2 SO 4 /0-5 N NaCl solutions at room temperature (Harston and Scully, 1969). Unstressed specimens exhibit intergranular corrosion but the application of a stress causes transgranular stress corrosion cracking. In that system the chloride ion appears to behave as an inhibitor. If the concentration is lowered to 0 1 N NaCl the corrosion rate increases and the fracture transition is not then seen: the stress corrosion is intergranular also, but the boundaries indicate incipient transgranular cracks that become wide blast fissures or elongated pits. In this case, increasing the overall corrosion rate removed the differences between the relative reactivities

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of the transgranular and intergranular paths. Typical fractographs are shown in Figs 10 and 11. Other austenitic stainless steels examined showed neither intergranular corrosion nor stress corrosion, a result that suggests that the alloy chemistry is of overriding importance. Other alloys also show preferential pre-existing intergranular paths. Titanium (Sanderson and Scully, 1968) and zirconium (Majumdar et ai, 1980) metals exhibit intergranular corrosion in unstressed specimens and intergranular stress corrosion cracking in CH 3 OH/HCl mixtures. When alloyed the higher strength alloys of both exhibit transgranular stress corrosion. Ferritic low strength steels exhibit intergranular corrosion in unstressed specimens anodically polarized in N H 4 N 0 3 solutions in which they crack intergranularly when stressed under open circuit conditions (Parkins, 1971). Attempts to

FIG. 10. A transgranular fracture of a Type 304 steel in a 5 N H 2 S 0 4 + 0-5 N NaCl solution obtained at room temperature (Harston and Scully, 1969).

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explain the cause of this increased chemical reactivity of the intergranular path concluded that it was due to unspecified substitutional solute elements rather than interstitial elements such as C or N (Flis and Scully, 1968). Since the presence of the interstitial elements is required for susceptibility to cracking an additional effect must be involved in the fracture process, e.g. a mechanical effect. Where the crack path follows a path that is not pre-existing the stress must act to create it and such paths are usually designated as strain induced, caused by plastic deformation. While the amount of

FIG. 11. An intergranular fracture of a Type 304 steel in a 5 N H 2 S 0 4 + 0 1 N NaCl solution obtained at room temperature. This situation caused a higher corrosion rate than that used to obtain Fig. 10. Wide fissures into the grains can be seen, which, under less corrosive conditions, would have developed into transgranular cracks. (Harston and Scully, 1969.)

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work done on single crystals is not large, that which has been done has shown consistently that stress corrosion cracking is observed only at and above the yield point, corresponding to the onset of plastic deformation (Forty, 1961). The strain-generated path is therefore always associated with the chemical properties of dislocations. In polycrystalline specimens threshold stresses, which are not always easy to discern, are often below the 0-1 % Proof Stress, a result that reflects the fact that in such material yielding occurs at these low values in some grains. The strain-generated path on an atomic scale has variously been considered to stem from the rearrangement of atoms about a dislocation. The strain energy associated with a dislocation in a pure metal is insufficient to cause electrochemical reactivity significantly enhanced above that of the undislocated matrix (e.g. Tromans and Nutting, 1965), and the rearrangement is always considered to include some enrichment of solute atoms to the dislocation line or some impoverishment, to such an extent that on an atomic scale significant differences in chemical reactivity are generated. Such effects have been associated with the occurrence of stacking faults (Swann and Nutting, 1959/60; Swann, 1963) in susceptible face centred cubic alloys and with heterogeneities arising from short-range order (Swann and Pickering, 1963). The local compositional differences that may promote marked differences in reactivity may be accentuated, or indeed mainly caused, by repassivation differences. Additionally, such differences may accelerate de-alloying processes (Swann, 1971) where these are important, perhaps injecting vacancies, causing a highly localized loss of ductility (Forty, 1960). The local morphology of cracking on this fine scale is difficult to discern and requires high resolution electron microscopy, particularly using a high-voltage instrument (e.g. Seamans and Swann, 1978). The shape and direction of the dissolving front have been subjects of investigation, e.g. tunnel corrosion (Swann and Embury, 1965), de-alloying sponges (Swann, 1971). If dissolution is the principal cause of crack propagation then the maximum dissolution current density that is possible will determine the maximum crack propagation rate. In practice, fractures change from one consisting predominantly of stress corrosion facets, to one consisting predominantly of ductile fracture features, and eventually to a completely ductile fracture as the stress, stress intensity or crosshead speed is raised in the different types of test that are used. In this context the maximum crack propagation rate refers to a predominantly stress corrosion fracture. The relationship between the

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maximum rate, v, and the current density, /, can be derived very simply from Faraday's Laws: (3) where J i s the chemical equivalent of the dissolving metal which has a density, p, and F is the Faraday constant. Correlations between propagation rates and maximum current densities have been made for a number of alloys (Parkins, 1977). Difficulties arise in interpreting the electrochemical experiments that are employed to derive the maximum dissolution rate for any particular alloy. The calculated maximum current density, for example, is not the maximum instantaneous value but the average over the time interval between successive transients of the type shown in Fig. 7. It will only approach the maximum value as the time interval decreases. Models have been proposed, for titanium alloys, for example, in which / exceeds 1000 A cm" 2 , a value higher than that observed under the relatively violent conditions of electrochemical machining (Beck, 1971). This is equivalent to a crack propagation rate in excess of 1 cm m i n - 1 , observed in titanium alloys, which is high, since in many alloys rates are well below that value, often closer to 1 mm h" 1 . Thus while equation (3) must be obeyed if an active path mechanism is operative, it has proved difficult to find an example where it is not obeyed. What would initially appear to be necessary, and perhaps unusual, appears to be almost unremarkable. Bare surface current densities on titanium alloys, for example, are the same for both susceptible and nonsusceptible grades (Beck, 1971). 2. Hydrogen embrittlement stress corrosion cracking Most susceptible alloys generate hydrogen at the crack tip and the possibility that a proportion of the hydrogen ions discharged enters the metal lattice at the crack tip and causes localized embrittlement must be considered. It can be difficult to establish that such a process is occurring but confirmatory evidence is found in the accelerating effect of adding cathodic poisons (e.g. Green et ai, 1976) of susceptibility increasing with the development of heat treatments that also increase susceptibility to hydrogen embrittlement (Thompson and Bernstein, 1980), in the effect of changes in loading mode upon susceptibility (St John and Gerberich, 1973), in measurements of hydrogen uptake (e.g. Gray, 1969) and in reversible embrittlement experiments (e.g. Scully and Adepoju, 1977) and other associated effects.

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It has proved very difficult to distinguish between the two possible causes of stress corrosion crack propagation since it has not been easy to establish a single, unambiguous experiment which enables a clear distinction to be made between them. Instead, an accumulation of data is made which provides evidence for one in preference to the other. Even here caution is necessary since there is no reason that both general mechanisms should not contribute to a fracture process. Under open circuit conditions the local crack tip will include both anodic and cathodic sites. The point at issue is the relative importance of each to the overall fracture process and how this can be discerned. While in some alloy systems one such reaction is of overriding importance, in others the contribution of each may be of comparable importance. In the alloy systems discussed briefly below the debate about active path corrosion and hydrogen embrittlement mechanisms is continually present, and possible explanations of any experimental data are sought with respect to the two possible explanations. In many cases it is not possible to provide an unequivocal explanation for any particular effect. Metallurgical changes, for example, the effect of variations of either major or minor alloying constituents, and electrochemical effects, for example, the effect of changes in potential, do not usually lead to completely clear explanations, particularly when associated film formation effects are also considered. IV. Alloy systems A. Titanium

alloys

Titanium alloys have a long history of laboratory-observed stress corrosion failures. A list of the main cracking environments is given in Table 1 and the subject has been extensively reviewed (Feeney and Blackburn, 1971). In practice very few failures have been reported, an apparent discrepancy that can be accounted for by a combination of initiation difficulties, high threshold values and unencountered environments. These points are of considerable importance, particularly the first two, since they determine the probability of cracking and that, for many purposes, is more important than the possibility. With so many different types of environments able to cause cracking there has been much discussion about the number of mechanisms that may be operative among the many different failures. One common thread running through many of these, although not all,

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is the possible role of hydrogen, since there is a long record of the adverse effects of hydrogen upon the mechanical properties of titanium and titanium alloys (Williams and Jaffee, 1958; and Williams, 1962). Two general effects were discerned. First, a high strain-rate hydrogen embrittlement was observed associated with a lattice containing sufficient hydrogen to ensure that the hydride phase was present. The consequent effect was a low impact strength with the region adjacent to the hydride, or the hydride phase itself, providing a large number of nucleating sites for cracks under conditions of high loading rates. The second effect was a low strain-rate hydrogen embrittlement which was found in hydride-free lattices supersaturated with hydrogen, which, under conditions of low loading rates, such as are encountered in a conventional tensile test, formed a plastically induced hydride phase that resulted in a reduced ductility that caused local embrittlement, including a smaller, more shallow type of dimple (Haynes and Maddocks, 1969). The observation that high strength α-titanium alloys were susceptible to cracking in sea water was reported by Brown (1966). Fracture was subsequently discovered to occur by a transgranular

FIG. 12. A transgranular fracture of a Ti-O alloy obtained in a CH 3 OH/HCl solution. The fracture consists of a mixture of cleavage and fluting failure modes. (Scully and Adepoju, 1977.)

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cleavage mode on a plane 14-16° from the basal plane (Blackburn and Williams, 1969). A typical fracture is shown in Fig. 12 (Scully and Adepoju, 1977) which illustrates a mixture of cleavage and associated fluted fracture (Aitchison and Cox, 1972) which is a low-energy ductile failure mode. The crack propagation rate resulting in such fractures in aqueous and organic media at room temperature can exceed 1 cm m i n - 1 . Similar fractographs occur as a result of liquid metal embrittlement in mercury, where rates can exceed 300 cm m i n - 1 , and also in molten salts (Feeney and Blackburn, 1971). Earlier work (Mori et ai, 1966) had reported the intergranular cracking of titanium in CH 3 OH/HCl and CH 3 OH/H 2 S0 4 mixtures, a phenomenon that could be inhibited by additions of small amounts of water ( < 5 wt. %). Subsequently it was observed that identical cleavage to that observed in aqueous solutions could occur in the higher strength α-titanium alloys as well as the intergranular mode (Sanderson and Scully, 1968). In the commercial Ti-6Al-4V and Ti-8Al-1 Mo-1V a + ß alloys only the a phase exhibits cleavage. The BCC ß phase is immune to stress corrosion cracking although this is a compositional not a structural effect since other /?-titanium phases exhibit (100) cleavage in the same environments (Feeney and Blackburn, 1971). A hydrogen mechanism has been proposed for this fracture also (Lycett and Scully, 1979). In the a phase the principal elements promoting cleavage are Al and O although this has not been systematically investigated. In 1967-8 a relatively simple mechanism was put forward (Sanderson and Scully, 1968) to explain these phenomena, based mainly upon investigations into a alloys. In the corrosive CH3OH/HCI mixture intergranular cracking was the result of an active path mechanism associated with the presence of some unspecified segregated species in solid solution in the grain boundaries. In unstressed specimens, for example, intergranular corrosion occurred under both open circuit and anodic polarization conditions. The application of a stress causes a more rapid penetration of the grain boundaries. In the CH 3 OH/HCl mixture and in the non-corrosive aqueous NaCl solution it was proposed that transgranular cleavage was caused by the hydrogen absorbed at the crack tip. It was argued that the interaction of the flux of absorbed hydrogen and the volume of metal immediately ahead of the tip caused à highly localized slow strain-rate hydrogen embrittlement effect (Powell and Scully, 1968). A hydride nucleus was formed which raised the shear stress locally and promoted cleavage. It was demonstrated that a hydride was formed when the passive film was broken in an acidified NaCl solution of

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pH 1 but not of pH 2, whereas no hydride was observed in N a 2 S 0 4 solutions of pH 1 (Sanderson and Scully, 1966). In neutral sea water titanium is passive and hydrogen absorption to any significant level occurs only while the film is broken since the film is highly impermeable to hydrogen. Considerable emphasis was placed therefore upon the repassivation time since this controlled the rate and thereby the amount of hydrogen absorbed. The results shown in Fig. 9 have already been discussed and were obtained from work in this system. The protective effect of cathodic polarization was also explained as the result of an accelerated repassivation rate, a point that was consistent with the observation that in 11 N HC1 solutions cathodic polarization had no effect since in this solution film formation would not occur (Powell and Scully, 1968). The schematic drawing in Fig. 13 attempts to relate the rate of hydrogen pickup to the effect of the imposed crosshead speed in raising the stress and to the lowering of the cleavage stress by absorbed hydrogen. The essential condition is that stress corrosion in neutral solutions occurs only when the repassivation rate is slow

Cleavage stress

Cleavage stress < ductile fracture stress when sufficient hydrogen is absorbed Intermediate strain-rate Low strain rate

Ductile fracture stress

r at high strain-rate

Time

FIG. 13. Schematic diagram showing the conditions under which a susceptible a-Ti alloy undergoes stress corrosion in slow strain-rate tests in neutral aqueous chloride solutions (cf. Fig. 9). At a high strain-rate the ductile stress is reached before sufficient hydrogen has been absorbed to lower the cleavage stress. At low strain-rates continued repassivation prevents a significant amount of hydrogen to be absorbed. At an intermediate strain-rate sufficient hydrogen is absorbed to promote cleavage fracture before the stress becomes high enough to cause ductile failure. Three kinetic processes are involved: the increase in stress with time, the absorption of hydrogen, and the repassivation process.

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STRESS CORROSION CRACKING

enough to allow a significant amount of hydrogen pickup while the stress at the crack tip is rising to a sufficiently high value in the same time interval. This is a relatively simple description but it emphasizes the specific dynamic interaction of several kinetic processes which must occur in a particular sequence and at certain rates before stress corrosion conditions are achieved. The proposed mechanism has subsequently had much support. It has been possible, for example, to demonstrate a removable embrittlement phenomenon in alloys exposed to environments from which hydrogen pickup is possible (Scully and Powell, 1970). Alloys exposed in the unstressed state in a CH 3 OH/HCl, for example, undergo intergranular dissolution, as already described but, in addition, hydrogen absorption occurs simultaneously (Menzies and Averill, 1968). This is drawn schematically in Fig. 14. Upon removal from the solution the specimens exhibit a three-zone fracture if broken

Grain boundary

Corroded grain boundary

FIG. 14. Schematic diagram showing that hydrogen absorption occurs together with intergranular corrosion when unstressed α-Ti alloys are exposed to a CH3OH/HCl solution. When such specimens are broken in air immediately upon removal from the solution in order to minimize the dispersion of the absorbed hydrogen cleavage is observed in front of the intergranular fracture. This zone of cleavage is not observed if hydrogen is allowed to disperse before fracture of the specimen.

immediately, consisting of intergranular and transgranular stress corrosion-like fractures and a dimple region. An example is shown in Fig. 15. If a time delay is inserted between the removal from the solution and the fracture then the intermediate transgranular fracture gradually diminishes in extent and when the delay is sufficiently great the zone is no longer seen. This gradual alteration is associated with an increase in ductility as shown in Fig. 16. A two-zone fracture is shown in Fig. 17. The effect of the time delay is to allow the absorbed

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FIG. 15. A three zone fracture obtained in a Ti-O specimen fractured in air immediately after exposure in the unstressed condition to a CH 3 OH/HCl solution. Intergranular cleavage and fluting, and dimple zones can be observed. (Scully and Adepoju, 1977.)

4 6 Time (days)

8

10

FIG. 16. The change in elongation of Ti-O specimens after removal from a CH3OH/HCI solution as a function of time lapsed between removal and fracture (Scully and Adepoju, 1977).

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FIG. 17. A two zone fracture obtained in a Ti-O specimen fractured in air 20 days after exposure in the unstressed condition to a CH3OH/HCl solution. Intergranular and dimple zones can be observed, unlike Fig. 15 where an intermediate cleavage zone was observed. During the 20 day period the absorbed hydrogen dispersed and was no longer available to cause cleavage. (Scully and Adepoju, 1977.)

hydrogen to diffuse away from the region immediately in front of the intergranular dissolution part drawn in Fig. 14. What these experiments show is that absorbed hydrogen can cause identical cleavage to that observed during stress corrosion crack propagation. The drawing of Fig. 14 assumes that grain boundary dissolution occurs at an inward rate that is slower than the diffusion rate of hydrogen, as a result of which difference a volume of metal ahead of the dissolving front becomes enriched in hydrogen. Under open circuit conditions this appears to be a realistic description. When anodic polarization is applied the grain boundary dissolution rate is increased whereas the diffusion rate of absorbed hydrogen in titanium is unaffected. The results drawn in Fig. 18 show that for a series of specimens galvanostatically polarized anodically the elongation-tofailure increases and then decreases with increasing current density. The transition point corresponds to an observed change from a threezone to a two-zone fracture when specimens are broken immediately after fracture (Scully and Adepoju, 1977). The interpretation of these results is that the current density at the transition point corresponds to the inward rate of grain boundary dissolution being equal to the hydrogen diffusion rate, with the result that an embrittled region

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ahead of the grain boundary dissolution front is not created. A cleavage zone is not therefore observed in the subsequent fracture process. Loading mode experiments of Green et al. (1976) have been interpreted as demonstrating the crucial role of hydrogen in causing transgranular cleavage in aqueous chloride solutions. Results are shown in Fig. 19. Susceptibility was marked in specimens loaded in mode I. Under such conditions hydrogen is considered to accumulate in regions exposed to high hydrostatic stress such as occur on the volume of metal immediately in front of the pre-crack. Under mode III loading conditions no such regions exist and, consistent with this, no susceptibility was observed. If an active path mechanism was the cause of fracture then crack propagation would be expected to occur at the same value of stress intensity factor independently of the loading mode since the requirement would be the creation of an active path at the crack tip as a result of film fracture and the stress condition

*f = * Δ - ( * α + € Η )

3 zone fracture

2 zone fracture

Applied current density

FIG. 18. The elongation-to-fracture of unstressed tensile specimens anodically polarized at various current densities in CH 3 OH/HCl for 2 h followed by immediate fracture in air. At low current densities 3 zone fractures were observed. At higher current densities 2 zone fractures were observed. The changeover is thought to correspond to the point at which the inward penetration of the grain boundaries is comparable with the inward diffusion rate of hydrogen. (Scully and Adepoju, 1977.)

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ahead of the crack would be of no consequence. Green et al. demonstrated that α-brass in ammoniacal solutions was equally susceptible under modes I and III as an example of such an active path mechanism. Green et al. also showed that additions of a cathodic poison exacerbated cracking, an effect to be expected if hydrogen is causing fracture and if, in the absence of the poison, less than 100 % of the discharged hydrogen is being absorbed. They observed also that the poison had no effect at the open circuit potential of -800mV(SCE) but had a marked effect at -500mV(SCE), an observation that was explained by considering the state of the poison at the two potentials. At — 500mV(SCE) As can be expected to deposit on the alloy surface from the solution and thereby promote hydrogen entry, while at — 800 mV(SCE) As will be present as the gas AsH 3 and therefore have no effect on surface reactions. Various fractographic and metallographic experiments have lent support to the hydrogen mechanism. It has been shown, for example, that finely divided hydride precipitates can cause cleavage in a susceptible α-titanium alloy fractured in air (Mauney and Starke,

Torsional loading

if ÎS

3-5 NaCl 3-5 NaCl + lOppm As

0-4 Time ( h )

10

100

FIG. 19. The effect of loading mode upon the fracture of Ti-8Al-lMo-lV specimens at -500mV(SCE) in 3 % NaCl solutions. Under torsion loading (mode III) stress corrosion fracture did not occur while under tensile loading conditions (mode I) stress corrosion was observed and was accelerated by the presence in the solution of arsenic. (Green et al, 1976.)

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1969). More recently, Koch et al. (1978) have concluded that the cleavage fractures in Ti-8Al-lMo-lV in 3 % NaCl solution and slow strain-rate hydrogen embrittlement are identical. The role of aluminium additions in promoting cleavage has not been clarified but Paton and Spurling (1976) have shown that the predominant {10Ï0} habit plane of the hydride in pure titanium is changed to basal and near-basal habit planes with increasing amounts of alloyed aluminium (3-6-6 wt. %). These observations are consistent with the possibility of stress-induced or strain-induced hydride formation as a factor in stress corrosion crack propagation since such additions promote transgranular cleavage ( ^ 5 % , Sanderson and Scully, 1968). Aluminium additions also promote slow strain-rate hydrogen embrittlement. A quantitative model of hydrogen embrittlement is still lacking. The highest crack propagation rates ( > l c m min" 1 ) are several orders of magnitude higher than the diffusivity of hydrogen, obtained by extrapolating high temperature data to room temperature (Powell and Scully, 1979b). A number of possible reasons can be put forward to account for such discrepancies. First, the extrapolation may give too low a figure. Diffusion rates along grain boundaries or dislocation pile-ups may not be reduced to such a low value as the matrix alone. Secondly, the diffusion rate of hydrogen in a volume of metal undergoing a low strain-rate may be higher than for a static lattice. Thirdly, the effect of absorbed hydrogen may be to initiate cleavage which then propagates over comparatively long distances through material that contains only residual level amounts of hydrogen. Under such circumstances, referred to as a long range effect (Powell and Scully, 1969a), the crack propagation rate can be much greater than the hydrogen diffusion rate since a small amount of hydrogen-induced embrittlement causes a large amount of cleavage and fluting fracture. Cracking in aqueous and methanolic solutions has been widely studied. For many of the organic solvents shown in Table 1 such studies have not been made. In many cases, though not necessarily all, cracking may be attributable to the residual moisture content, a feature that has not been demonstrated unequivocally. In N 2 0 4 cracking is observed (Sedriks et ai, 1969). It occurs with the formation of a thick black film. Sedriks et al. (1969) suggested that propagation is associated with its continual fracture and repair. In liquid metal embrittlement crack tip interactions of mercury and titanium atoms have long been advocated as causing fracture which can reach very high rates as has already been described. In both these

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environments hydrogen may not have a function in the fracture process. The transgranular cleavage observed in aqueous, methanolic, mercury and molten chloride environments is, therefore, probably caused by more than one process. While these may all involve lowering of cleavage stresses by weakening of bonds and be mechanical in origin, there are arguments for a chemical model in aqueous solutions. The Mass Transport Kinetic model of Beck (1971), for example, explains cracking in aqueous solutions as an active path process. The role of the chloride ion (and other halide ions) has not been fully explored. In aqueous solutions hydrolysis results in a low crack tip solution pH of c. 1 -7 (Brown et ai, 1969). Chloride ions also retard the repassivation process. Beck (1971) and his colleagues (Alkire et al.9 1978) have provided arguments and data on the formation of salt layers on the metal surface prior to repassivation. A mathematical model indicates that in a 3 N HC1 solution a salt film 2-10 nm thick forms and disappears within the time period 10 _ 5 -10~ 3 s after generation of oxide-free surface, during which time the anodic dissolution rate of titanium would be expected to exceed 150 A cm" 2 . Little is known about these layers prior to repassivation. The region II propagation rate is dependent upon the halide ion concentration and such effects have been interpreted as arising from the state and composition of the surface film (Scully, 1973). Increasing halide ion concentrations may cause a halide layer to exist for a longer time or, on average, cause contamination of the film for a longer time. It has also been suggested that hydrolysis of the surface halide layer may lead directly to hydrogen absorption, possibly aided by the adsorbed chloride ion (Scully and Adepoju, 1977). This subject is clearly of considerable importance but little is known about the nature of the surface and the function of the halide ion. The hot salt cracking process was investigated over 15 years ago by a number of laboratories (Raring, 1966) and in the intervening years not much additional work has been reported. The important role of moisture has been demonstrated but whether, for example, cracking can occur in the absence of moisture has not been shown. Mechanisms involving the formation of HC1 have been put forward but whether such formation is a cause of cracking or a result of other reactions causing cracking is not always clear. It has been demonstrated, however, that the fracture surfaces of hot titanium alloy specimens coated with solid NaCl absorb hydrogen very readily from atmospheres with extremely low relative humidities (dewpoint — 50°C) to levels exceeding 104 ppm (Gray, 1969). Ondrejcin (1970) showed

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clearly that hydrogen produced during the corrosion of titanium alloys by halide salts and subsequently absorbed by the titanium is responsible for their cracking. Fracture modes in hot salt cracking and high pressure hydrogen gas were similar. Stressed Ti-Al alloys fractured when bombarded with low energy protons, a result that emphasized the effect of hydrogen in causing fracture. B. Zirconium

alloys

Zirconium alloys have relatively limited uses and stress corrosion failures have not been very common in practice although a range of environments can cause cracking in laboratories, as reviewed by Cox (1975). In CH3OH/HCI mixtures the same two forms of cracking are seen as are seen in α-titanium alloys. Intergranular cracking is considered to be a stress-induced (or strain-induced) active path fracture mode. Intergranular corrosion occurs in the absence of stress but at a lower inward rate. This difference is illustrated in Fig. 20. The

200

< 40

00

<-?f

Time (h)

Stressed specimens

4

%

Q. T3

>v

σ 5 0

Unstressed specimens

1

0Ό0Ι

0ΌΙ Crosshead speed (^.m s"1)

0-1

FIG. 20. The rate of grain boundary penetration observed in Zircaloy-2 specimens exposed (i) in slow strain-rate experiments, and (ii) in the unstressed condition, in a CH3OH/HCI mixture. The difference between the two represents intergranular stress corrosion cracking. At very low strain-rates this difference disappears since over a long time sufficient hydrogen is absorbed to promote transgranular stress corrosion by cleavage before initiating intergranular stress corrosion. (Majumdar et ai, 1980.)

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139

figure shows the depth of intergranular separation observed in two types of specimens: unstressed, and strained over a range of constant crosshead speeds and both over the same time of testing. The difference in the depth of intergranular fracture represents the Stressor strain-assisted component. At the lowest value of crosshead speed there is no difference between the specimens tested unstressed and the specimens exposed to that crosshead speed. The explanation for this has been that during the relatively long exposure of specimens a lot of hydrogen is absorbed and by the time that the stress on the low crosshead specimens begins to rise to values able to cause fracture, transgranular cleavage is caused rather than intergranular cracking. No intergranular separation other than that caused by corrosion is therefore observed and both types of specimens therefore exhibit identical amounts (Majumdar et al). The cleavage observed is shown in Fig. 21 and is thought to be caused by absorbed hydrogen, although this view has been put forward only by the present author and his colleagues. Consistent with this view is the observation of a reversible embrittlement phenomenon (Majumdar and Scully, 1979) and the effect of a cathodic poison, Se0 2 , in promoting additional cleavage (Golozar and Scully, 1982). Figure 22 shows the types of fracture observed under open circuit conditions as a function of crosshead speed. The addition of the poison increases the highest value of crosshead speed

FIG. 21. A transgranular stress corrosion fracture of a Zircaloy-2 specimen obtained in a slow strain-rate test in a CH 3 OH/HCl mixture.

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at which cleavage is observed by over an order of magnitude. Arguments similar to those presented in Fig. 13 would explain this result. Cleavage results when sufficient hydrogen is absorbed before the stress reaches the ductile overlead value. Because the poison increases the absorbed flux, the upper crosshead limit is raised. Under conditions of long time exposure to a CH 3 OH/HCl mixture hydrogen blisters form in the grain boundaries. An example is shown in Fig. 23 (Majumdar et ai, 1980). This is a rather surprising observation since the dissolution of hydrogen in zirconium would be the expected preferential reaction. It must be supposed that the recombination of hydrogen atoms is catalysed by at least one of the intergranular precipitates. Figure 24 shows an inside section of a blister in which the intergranular aspect is clear. In aqueous chloride solutions cracking occurs only when anodic polarization is applied which raises the potential to the value at which film breakdown occurs. The type of fracture observed is transgranular and is partly covered with a corrosion product. In CH 3 OH/HCl

Transgranular >v stress corrosion tracking \ ^

Ductile failure

G I0 2 σ> c o σ u

B. 10 1 Γ

Q

I1 0-001

Intergranular stress corrosion cracking

1 L_ 0-01 0-1 Crosshead speed (/im s"1)

I

FIG. 22. The types of stress corrosion fracture observed in Zircaloy-2 specimens subjected to slow strain-rate tests in a CH 3 OH/HCl solution under open circuit conditions. The addition of the cathodic poison, Se0 2 , promotes the occurrence of the transgranular cleavage mode.

FIG. 23. Blister formation on a Zircaloy-2 specimen exposed in the unstressed condition to a CH3OH/HCl solution.

FIG. 24. Part of the inside surface of a blister illustrating the intergranular location of blisters.

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mixtures cracking is promoted by anodic polarization and retarded by cathodic polarization, and under open circuit conditions it is inhibited by additions of water. These results are similar to those observed in titanium alloys. Zircaloy-2 exhibits transgranular cracking in I 2 vapour at 350°C and has been a possible source of failure in cans in nuclear power stations. The mechanism of such failures has not been elucidated. C. Austenitic stainless

steels

Austenitic stainless steels exhibit stress corrosion cracking in hot chloride solutions and there is a long history of such failures in chemical plant and other establishments (Theus and Staehle, 1977). In aqueous solutions of concentrated chlorides cracking is commonly transgranular but in high-temperature pressurized water and steam intergranular fracture predominates. For historical reasons much laboratory work has been done in MgCl2 solutions boiling at c. 154°C. A film exists on the surface in this type of solution and cracking is associated with its localized rupture. Maximum propagation rates are in the range 1-5 mm h~ l and much data have been interpreted as being indicative of an active path mechanism. Early work attempted to demonstrate that, at the potential observed at the region of the crack tip, the dissolution current density was sufficient to be able to account quantitatively for such rates in accordance with equation (3). Under conditions of rapid solution movement reducing concentration polarization to a negligible level it was shown that such current densities could exist (Hoar and Scully, 1964). In practice it might be supposed that the continuous evolution of hydrogen at the crack tip would provide such a stirring effect, in addition to the continuous strain-rate and movement of the crack tip surface. Since there is no obvious transgranular pre-existing path for dissolution a strain-induced model is commonly invoked. Based upon various types of experimental observations by transmission electron microscopy dissolution is thought to occur on stacking faults or regions adjacent to them or other dislocation pile-ups from narrow trenches and possibly in the form of small tunnel formation at the dissolving front (Swann, 1963). Fracture occurs on {100} planes in a high nickel austenite (Reed and Paxton, 1961) or {210) or {110} in Type 304 austenite (Harston and Scully, 1970; Marek and Hochman, 1971). No explanation has been put forward for these fractures except possibly by Seamans and Swann (1978) who examined dissolution

STRESS CORROSION CRACKING

143

effects in stress corrosion specimens by high voltage electron microscopy. A typical result is shown in Fig. 25. Dissolution has occurred on {110} planes in <111> direction. The authors suggested that this may arise from a process of preferential repassivation or because {110} planes are the least reactive when unfilmed. These different aspects may, of course, be different manifestations of the same surface property, as is discussed below.

FIG. 25. Slot-like attack observed in miniature U-bends of Type 302 steel exposed to MgCl2 solutions (Seamans and Swann, 1978).

The role of hydrogen in promoting transgranular fracture has been emphasized by numerous workers over many years; others, whose work has been summarized by Staehle (1971), concluded that it has no role. Solutions of LiCl containing Li 2 S, for example, do not promote cracking or increase hydrogen permeation rates more than Li2S-free solutions. Absorbed hydrogen promotes the formation of ε and a' martensite in some austenite lattices (Holzworth and Louthan, 1968). The latter phase has been observed on austenite stress corrosion fracture surfaces (Birley and Tromans, 1971). It has been argued that such a transformation is not an essential part of the fracture mechanism since stress corrosion occurs above the Ms temperature of some austenites. This latter point has been investigated recently by Liu et al. (1980). In MgCl2 solutions boiling at 154°C a Type 310 steel was found to crack on or near to {100} planes. Fractures in Type 304 steel were more complex and fracture orientation was scattered. It was not {100} but fell into two distributions, one near to {211} and the other near to {110}. Both a' and v. martensites were determined by electron

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J. C. SCULLY

diffraction on Type 304 fracture surfaces but these could not be found on 310 steels. At the higher temperature of 289°C no martensite phases were detected on Type 304 steel. The fracture facets were large and flat, similar to those observed on Type 310 steel. The facets obtained from specimens broken in gaseous hydrogen or after cathodic charging at room temperature were too small for precise orientation determination. The results are consistent with a' martensite forming in tests done at temperatures below the Ms and with cracking occurring both below and above Ms. Eliezer et ai (1979) had concluded previously that austenite can be embrittled and can suffer stress corrosion without the presence of martensite. According to Bursle and Pugh (1979) cracking under normal constant load conditions proceeds by repeated cycles of embrittlement of the lattice ahead of the crack tip with limited cleavage fracture through the region, the crack advancing a distance of Ax per event with a time interval At between events. Hahn and Pugh (1980) applied a load pulsing technique to studying propagation in a Type 310 steel in boiling MgCl 2 . By relating the observed fractographic marking to the timing of the pulses the results indicated Ax to be c. 0-5 μπ\ and Δ/ to be c. 15 s. The metallographic observations of Hanninen and Hakkarainen (1979) have lent some support to the argument that the transgranular crack is propagating as a result of embrittlement induced by absorbed hydrogen. They charged a Type 316 steel cathodically at 80°C in a 1 N H 2 S 0 4 solution containing 250 ppm of NaAs0 2 for 18 h at 50 mA cm" 2 . Specimens broken immediately exhibited no ductility and gave rise to a variety of transgranular fractures characteristic of those reported after stress corrosion failure in boiling MgCl2 solutions. After 5 days at 100°C, charged specimens exhibited considerable ductility and ductile fracture, apart from some shallow surface cracks, results that are characteristic of an outgassing phenomenon. Such results are similar to those reported for α-titanium alloys and described above. For austenitic stainless steels it appears that hydrogen can cause the characteristic flat fractures commonly observed in stress corrosion cracking, two different examples of which are shown in Figs 26 and 27 (Harston and Scully, 1970). Whether dissolution can cause the same fractures is less clear, despite the observed faradaic equivalence, and is perhaps less easily demonstrated. Most observed effects do not provide an unequivocal answer, e.g. the effects of alloying elements. The influence of a number of alloying elements is complicated since there are interactions between individual elements (Jones and Hines,

STRESS CORROSION CRACKING

145

1961) and, also, because they have several effects not all of which may contribute to susceptibility to the same degree. Nickel has long been recognized as beneficial and phosphorus as detrimental, as well as platinum (Latanision and Staehle, 1969). While lowering the C and N content below the normal commercial levels was shown long ago to have a beneficial effect, interpretation is complicated by such changes causing a structural alteration since such steels are then BCC. While BCC stainless steels are not commonly susceptible to cracking in MgCl2 solutions which might explain the C and N effect, caution must be exercised since it has been shown (Bond and Dundas, 1978) that small alloying additions of Ni and/or Cu can render BCC stainless steels susceptible to transgranular cracking and result in fractographs that appear to be very similar to those observed in austenitic steels.

FIG. 26. Transgranular fracture of a Type 304 steel after exposure to a MgCl2 solution at 154°C (Harston and Scully, 1970).

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J. C. SCULLY

Crack propagation rate measurements in MgCl2 solutions as a function of crosshead speed show a lower constant range followed by an increasing range up to 100% ductile fracture as shown in Fig. 28 (Talebian and Scully, 1981). These values are determined by the potential and the solution composition and temperature (Takano, 1974). Cathodic polarization lowers and eventually prevents crack propagation because of a repassivation effect (Robinson and Scully, 1978). In addition, the predominantly transgranular fracture changes to intergranular as the potential is lowered and as the temperature is reduced. In general, amounts of intergranular cracking have been variously reported in MgCl2 solutions. A recent study has shown that the fracture path in non-sensitized Type 304 steels is determined by

FIG. 27. Transgranular fracture of an 18Cr-18Ni austenitic steel after exposure to MgCl2 at 154°C (Harston and Scully, 1970).

STRESS CORROSION CRACKING

147

the composition of the liquid at the crack tip (Talebian and Scully, 1981). This has been hypothesized as having a range of compositions which, in turn, cause a range of repassivation rates. In the less aggressive solution in which repassivation is more rapid cracking is intergranular whereas in the more aggressive solution in which repassivation is less rapid cracking is transgranular. Inhibitors, cathodic polarization, lower temperatures and Mo additions all promote increasing amounts of intergranular cracking. Such an analysis emphasizes the primary importance of the crack tip chemistry in determining crack tip path and attributes no direct role to any mechanical variable in such determination. This subject is discussed below in section V.

500

>

300

Crosshead speed (nm s ) FIG. 28. The relationship between stress corrosion crack velocity of Type 304 steel in MgCl 2 solutions at 154°C and crosshead speed (Talebian and Scully, 1981).

D. Alpha brass Although a wide range of copper alloys exhibits stress corrosion cracking most attention has been focused upon commercial 70Cu30Zn type alloys exposed to ammoniacal environments. While it is an old form of failure, examples still occur, in condenser tubes, for example. In laboratory work the role of pH, potential and Cu 2 + /NH4 composition has been carefully formulated by Mattsson (1961) and Johnson and Leja (1966).

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J. C. SCULLY

In neutral solutions cracking tends to be intergranular under conditions of a high rate of tarnish film formation. If the tarnish rate is lowered cracking becomes transgranular (Kermani and Scully, 1979a), a result reported also for the imposition of cold work (Pugh, 1971). This effect is identical to that described above for austenitic stainless steels. The exact importance attached to the tarnish film during crack propagation has rather diminished in recent years. Whereas its formation and fracture were thought to be an important part of the crack propagation process, perhaps even the actual stage (Forty and Humble, 1963), the work of Bertocci (1969) revealed that it forms by a precipitation process on top of a thin underlying film which affords the protection more closely associated with a passive film in transition metals (Gabel et ai, 1976). The exact relationship between the conditions of formation of the two has not been determined and most descriptions tend to focus upon the outer, thicker film. The relationship between propagation rate and crosshead speed is shown in Fig. 29. The values are dependent upon the solution

1

E

I02

,E

*— ' >N o o ω

> I01 _*: o o \_ o c o

<> / o I—

Ü

o

10°

(Λ Φ

ω VLO

I0"5

I0" 4 I0" 3 I0" 2 I0"1 1 Crosshead speed (cm min" ) FIG. 29. The relationship between stress corrosion crack velocity of 70Cu-30Zn brass in a neutral ammoniacal solution and crosshead speed (Kermani and Scully, 1979a).

149

STRESS CORROSION CRACKING

composition and specimen potential, as is the cracking morphology. The effect of pH changes is shown in Fig. 30 and of potential in Fig. 31. Anodic polarization causes a drop in propagation rate which has been explained as being due to the non-formation of the tarnish film as a result of which the same anodic current is distributed to a large number of sites. At high values of pH tarnish films form which are less protective. In unstressed specimens, for example, intergranular corrosion occurs (Pugh, 1971). Under imposed stress conditions cracking can be intergranular or transgranular. For the intergranular cracking observed in ammoniacal solutions there is much agreement that an active path mechanism is operative. Why this should occur has not been determined beyond supposing that some minor segregated impurity is able to cause high localized corrosion rates either by increasing dissolution rates and/or by preventing or delaying protective film formation. Crack propagation is continuous on a very fine scale as determined by acoustic emission

50l· 40 l· 1

JZ 30l· E

o

I

20 l· I0l· I

I 5

I 6

pH

i 7

i 8

I 9

FIG. 30. The effect of pH on the stress corrosion crack velocity of a 70Cu-30Zn brass at a crosshead speed of 0-33 μπι s" 1 (Kermani and Scully, 1979a).

150

J. C. SCULLY

and by pulsed load techniques similar to those employed in examining cracking in austenitic stainless steels. Transgranular cracking has caused more discussion. Susceptible FCC alloys have a low stacking fault energy and selective dissolution of such regions was proposed over twenty years ago (Swann and Nutting, 1959-60) and selective attack on dislocation pile-ups by ammoniacal vapour on thin foils of a-brass was observed by Tromans and Nutting (1965). A mechanism based on the dissolution of straininduced paths arose. This mechanism has been challenged by the work of Beavers (1977) and Tong (1977) who have examined in considerable detail the fractography of transgranular cracking.

Intergranular

12

o o > σ

c o o vo o ω ■*=

2

-40

-20

20

40

60

80

100

F(SCE) (mV) FIG. 31. The effect of potential upon stress corrosion crack velocity and fracture mode of a 70Cu-30Zn brass in a neutral ammoniacal solution at a crosshead speed of 0-33 μΐΐΐ8 _ 1 (Keramani and Scully, 1979a).

STRESS CORROSION CRACKING

151

Fracture follows {110} planes and single crystals give rise to matching interlocking fractures which resemble cleavage fractures. Pulsed loading experiments show that fracture occurs over relatively large distances in very short time intervals, a process consistent with an embrittling mechanism. What has not been discerned is what process might cause any embrittlement. Localized dezincification has been advocated (Forty, 1960) and while there is some evidence for it (Pinchback et ai, 1975) it has not been demonstrated as occurring in depth. Since hydrogen can cause embrittlement in many other alloys it must be considered for brass. The potential at which cracking occurs does not suggest that absorbed hydrogen could have any role in the fracture process because it is too noble and the possibility of a large iR drop down the crack can be discounted. There is, at the moment, no other obvious source of hydrogen atoms. A process of chemical cleavage has been suggested (Kermani and Scully, 1979b), i.e. the development of a rational crystal plane of separation as a result of a directional dissolution process, with the chemical attack creating a slot which grows mechanically before being arrested. The embrittling action would be the consequence of an unspecified anodic process, e.g. dezincification. Cracking has been analysed as occurring on {110} planes on "parallel but displaced planes" with the steps between the fracture facets also being crystallographic, occurring on {111} planes as a result of a shear process (Beavers, 1977) but {100} and {111} planes have also been identified (Kermani and Scully, 1979b, c). As with the transgranular fractures in austenitic stainless steels, the transgranular fractures in brass exhibit flat cleavage-like surfaces both in neutral and strongly alkaline solutions as shown in Figs 32 and 33. Whereas results on the steels have indicated possible hydrogen effects, including a reversible embrittlement phenomenon, such results have not been reported for brass. E. Aluminium

alloys

Aluminium alloys have a long history of intergranular failures in aqueous and organic environments containing traces of moisture. The subject has been extensively reviewed by Speidel (1972). While a film fracture process must be invoked as a necessary part of the crack propagation mechanism, the subsequent critical event at the crack tip that is responsible for fracture has not been clearly discerned. An active path mechanism was long advocated (Mears et al, 1945) with attention being given to the galvanic effects created by precipitate/matrix potential differences, some precipitates being anodic,

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some cathodic. Heat treatments giving maximum hardness and microstructures that strained by forming narrow bands of very high dislocation densities were related to high levels of susceptibility without at the same time providing unambiguous mechanistic insight.

FIG. 32. A transgranular fracture of a 70Cu-30Zn brass obtained in a neutral ammoniacal solution (Kermani and Scully, 1979c).

FIG. 33. A transgranular fracture of a 70Cu-30Zn brass obtained in a 15 N NH 3 solution (Kermani and Scully, 1979b).

STRESS CORROSION CRACKING

153

In recent years more attention has been given to the possible role of hydrogen. Evidence for its role has arisen from experiments illustrating reversible embrittlement (Gest and Troiano, 1972) and the liberation of hydrogen from embrittled alloys broken in vacuum (Montgrain and Swann, 1974). Recently the effects of strain-rate, hydrogen embrittlement and electrode potential have been delineated. For a 7049 alloy in the T651 condition cracking was transgranular at low levels of absorbed hydrogen but intergranular when that level rises (Holroyd and Hardie, 1981). Two regions of intergranular cracking were observed at potentials more anodic and more cathodic than the open circuit corrosion potential. A reversible embrittlement was observed. Slow straining in an inert environment was necessary for total recovery. Moisture in the environment may mask recovery by causing additional embrittlement. The major cause of embrittlement is hydrogen embrittlement. Different aluminium alloy systems may behave differently, as might be expected, but a picture is emerging. Absorbed hydrogen can cause intergranular cracking. For Al-Mg alloys, there appears to be sufficient current flowing to establish faradaic equivalence and an active path mechanism has been invoked (Ford, 1977). Crack propagation in some Al alloys is markedly irregular, even being accompanied by audible signals (Watkinson and Scully, 1970), is very much affected by changes in hardness, and in general appears to be mechanical in mode. This need not be true for all alloys, however, or at all propagation rates, and an active path contribution cannot be ruled out. Loading mode experiments of Green et al. (1976) suggest this conclusion. They observed that for alloy 7075-T6 susceptibility in mode I loading was much greater than for mode III loading but that some cracking did occur in mode III. This could be attributed to active path cracking. This result, when considered in conjunction with the work of Holyroyd and Hardie, suggests that cracking can occur both by hydrogen absorption and active path distribution, at least, in the alloys tested. V. Additional mechanistic features A. Crack

morphology

As can be seen in Table 1, it is commonly the practice to indicate the fracture path that is observed in typical failures. Such observations have two possible benefits. First, they are likely to be helpful in the

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analysis of service failures. Intergranular fractures of low strength steels, for example, are likely to be environmentally induced since there is no other obvious cause of such failures in these alloys. Secondly, the existence of a commonly observed crack path might be espected to form part of any proposed mechanism since the preferred path must be explained. What, perhaps, needs to be emphasized is that in most susceptible alloys the commonly observed fracture path is not the only one that can occur. While it is usual to see one particular kind—intergranular or transgranular—under service conditions or in laboratory experiments, with little exception, either fracture path is possible. In formulating mechanisms of cracking, therefore, the existence of possible alternative paths will clearly lead to a different emphasis being placed upon the significance of the common morphology since it is not unique, merely preferred. It is possible to offer a rationale for such observations and this is attempted, although the picture is not completely clear. Crack initiation occurs as a result of the corrosion of an alloy surface, arising from a highly localized film breakdown, as has been discussed at length in section III. The exact conditions of crack initiation will depend partly upon the type of loading arrangement and partly upon the environment that is employed. In Constant Extension Rate Test experiments or in specimens loaded or strained to produce yielding, for example, emergent slip steps will fracture the surface film and reveal bare metal as drawn in Fig. 1 unless this is of a thickness comparable to the emergent step size. Fracture of the film can be ascribed to a mechanical cause insofar as fresh metal is generated mechanically. In specimens under low loads or strains pitting may precede crack initiation, i.e. corrosion attack may not necessarily be influenced by the imposition of a stress. Such an event may help the crack to initiate either mechanically, by locally raising the stress, or, electrochemically, by causing localized solution changes that break down the film. Such latter events may, for example, result from anodic polarization. Such pitting may start at grain boundaries or at surface inclusions or at general surface irregularities. In many systems of intergranular cracking an active path mechanism is invoked, e.g. α-Ti alloys, Zircaloy-2, 70Cu-30Zn alloys, austenitic stainless steels and low strength Fe-C steels, and all of these have been cited in previous sections. Aluminium alloys could also be included in this test but, as indicated in the previous section, the point is disputed. In all the alloys indicated, apart from aluminium, acoustic emission studies and pulse loading techniques, for example, have both failed to provide evidence for the occurrence of irregular intermittent

STRESS CORROSION CRACKING

155

steps in the propagation process. By contrast, these are detectable in transgranular propagation. The question then arises as to whether a transition will occur. If an active path is the only operative mechanism for either mode of fracture then a transition will occur only if the incipient strain-induced transgranular path is chemically more reactive than the intergranular path. Various factors must be included in considering the origin of the apparent competition between the intergranular path being followed and the possible new transgranular path. For example, the relative rates of repassivation of intergranular surfaces and slip step surfaces may be markedly different. In addition, the thickness of the films forming on both types of surfaces may be important. Thick films on mild steel and 70Cu-30Zn surfaces, for example, are associated with the occurrence of intergranular fracture paths. Consideration of the event drawn schematically in Fig. 6 may provide at least a partial explanation. When thick films are broken it can be supposed that less fresh metal is revealed than is likely to be the situation with thin films for the reason drawn schematically in Fig. 34. The assumption is made that the film shears in the simplest possible way. The amount of fresh metal accessible to the solution is small compared with that arising from a comparable event on a metal surface covered with a thin film. If a hydrogen embrittlement mechanism is possible in addition to an active path mechanism then it is necessary for hydrogen to be absorbed and diffuse ahead of the dissolving crack tip to a sufficient concentration to initiate a hydrogen fracture as, for example, has already been described above for α-titanium alloys (Fig. 14). For these and for Zircaloy-2 the transgranular cleavage fracture occurs in preference to the intergranular path because it is more rapid and appears to need less associated corrosion. Transgranular cleavage occurs, for example, in neutral aqueous solutions in which intergranular corrosion is not observed because repassivation is too rapid. In the low-strength α-titanium grades cleavage does not occur so these alloys cannot fail in neutral aqueous solutions. The same analysis may apply to austenite: intergranular fracture from a dissolution mechanism and transgranular fracture, perhaps, from absorbed hydrogen, although this latter point remains to be proved, as has been indicated above. Unlike titanium and zirconium alloys, however, the transgranular fracture mechanism in austenitic stainless steels appears to require more associated corrosion, not less, than that required for intergranular cracking and is associated with a lower repassivation rate than that required for intergranular cracking. For 70Cu-30Zn brass a similar possible analysis can be attempted although no

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J. C. SCULLY

embrittlement mechanism has been demonstrated up to the present moment. The higher strength aluminium alloys exhibit intergranular cracking attributable to both an active path and a hydrogen embrittlement mechanism in a proportion that has not been determined. Transgranular cracking has been reported (Watkinson and Scully, 1970). It has been explained as occurring in a 7049 alloy at an early stage when the volume of metal in front of the crack tip contains hydrogen below the level required to cause intergranular cracking (Holroyd and Hardie, 1981). At a later stage the grain boundaries contain sufficient hydrogen and intergranular cracking then occurs. The descriptions given above are summarized rather simply in Fig. 35 which indicates the principal sequences of events surrounding crack initiation and the selection of a path. Film breakdown, either by mechanical or chemical means, is followed by intergranular corrosion Thick film

Thin film

FIG. 34. Schematic diagram indicating that the thickness of the film, /, affects the amount of fresh metal revealed by an emergent slip step.

157

STRESS CORROSION CRACKING

Intergranular corrosion

Intergranular

I cracking

B

Transgronular

Crack initiation after film breakdown

(cracking

Time-

FIG. 35. Schematic diagram which summarizes the type of fracture observed as a function of time. Examples of paths followed are given in Table 2.

TABLE 2 Stress corrosion fracture paths observed in various systems (drawn from the Schemes shown in Fig. 35) Path

System

A+ B A + B+ C B

Ti and Zr in CH 3 OH/HCl Ti and Zr alloys in CH 3 OH/HCl Mild steel in hot nitrates, hydroxides and carbonate/ bicarbonate solutions Al alloys in aqueous solutions α-Brass in neutral ammoniacal solutions Zr alloys in CH 3 OH/I 2 Austenitic stainless steels in high-temperature steam Austenitic stainless steels in hot concentrated chloride solutions High strength steels α-Brass in strongly alkaline ammoniacal solutions Ti and Zr alloys in aqueous solutions Mg alloys in aqueous solutions

B H- C or C

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J. C. SCULLY

or stress corrosion or transgranular stress corrosion. Intergranular corrosion may change to transgranular stress corrosion as described. Transgranular stress corrosion does not change to intergranular cracking unless the solution at the crack tip changes as is described later in this section. Even Fig. 35 is complete only under certain conditions concerning the crack tip chemistry. It has been argued that for austenitic stainless steels, 70Cu-30Zn brasses and mild steels that the fracture path is determined mainly by the composition of the solution immediately adjacent to the crack tip since this determines the repassivation rate (Scully, 1980). In relatively aggressive solutions, in which repassivation rates are relatively slow, cracking is transgranular, whereas in relatively less agressive solutions, in which repassivation rates are relatively rapid, cracking is intergranular. Typical conditions that give rise to these changes are given in Table 2. In austenitic stainless steels, for example, in Constant Extension Rate Test experiments, it is possible to produce 100% transgranular fractures or virtually 100% intergranular fractures in non-sensitized material, merely by controlling the potential at different values. The less noble potential, nearer to complete repassivation, gives intergranular cracking because the crack tip solution is less aggressive and repassivation is more rapid. In other cases, and at intermediate values of potential, the transition may not be so marked in going from one mode completely to the other but show definite transitions from one mode to a mixed mode. This would be an explanation for mixed mode fractures which could arise as a result of very small changes in solution composition occasioned by variations in localized conditions which may be caused in a number of ways, e.g. rapid increases in corrosion around inclusions etc. could create differences in repassivation rates as a result of the ensuing hydrolysis. In thin sheet specimens it has been argued that yawning of cracking specimens towards the end of Constant Extension Rate Test specimens causes partial changes in the crack tip solution as a result of the drawing in of bulk solution towards the crack tip (Talebian and Scully, 1981). Such marginal changes produce a final intergranular region in austenitic stainless steels, as is depicted in Fig. 36 which also shows an initial intergranular region. The type of reversion—transgranular to intei granular—is not depicted in Fig. 35. It has been seen at high values of crosshead speed (Talebian and Scully, 1981). The explanation of the results drawn schematically in Fig. 36 is depicted in Fig. 37 which attempts to show how the crack tip chemistry affects the fracture morphology in Constant Extension Rate Test experiments on notched specimens. Stress

STRESS CORROSION CRACKING

159

corrosion cracking occurs as the result of an interreaction of several kinetic reactions. In this example the rate of hydrolysis is considered in relation to the rate of stress increase. The composition of the solution at the tip of the notch gradually changes towards the composition to be found at the crack tip which causes transgranular cracking. At a prior stage, however, it is suggested that the solution becomes sufficiently aggressive to cause intergranular cracking but not transgranular, corresponding to the condition described as less Intergranular

~F

T~

N

o T

Transgranular

Ductile

c

H

FIG. 36. Schematic drawing of fracture surface of Type 304 specimens broken in a slow strain-rate test in MgCl2 solutions at 154°C illustrating three zones in which a predominant type of fracture is indicated.

FIG. 37. Schematic diagram of the sequence of events occurring at the tip of a notch in specimens exposed to a number of strain-rates. At high strain-rates the stress rises relatively rapidly and is able to initiate stress corrosion at a time when the solution hydrolysis is incomplete and can cause only intergranular cracking. At low strainrates by the time the stress has risen to the stress corrosion initiation value hydrolysis reactions have reached an equilibrium point and can cause transgranular cracking. (Talebian and Scully, 1981.)

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J. C SCULLY

aggressive. At the same time the stress is increasing. At high values of crosshead speed the fracture initiation stress is reached while the solution is in that condition and intergranular cracking occurs. At lower values of crosshead speed the solution is able to cause transgranular fracture by the time that the stress has reached the fracture initiation value. Figure 37 is intended to be a simplified representation of what will be rather more complicated interrelated events. For example, the initiation stress for the two forms of cracking may not be identical, and the rate of hydrolysis can be expected to be related to the imposed crosshead speed since this determines the dissolution rate. Nevertheless, in its simple form it explains the occurrence of the observed range of fractures and serves to emphasize in this and the other two systems of stress corrosion shown in Table 2 that the fracture morphology is determined by the crack tip chemistry alone and not directly by any mechanical variable such as stress or strain-rate. For example, it has been shown that if the bulk solution is made of a composition corresponding to the more aggressive kind that causes transgranular cracking then the stress corrosion part of the fracture is 100% transgranular even at high values of crosshead speed (Talebian and Scully, 1981). The results obtained on austenitic stainless steels in boiling MgCl2 solutions may differ from those obtained in other environments in which the repassivation differences are either less or greater. At high temperatures, e.g. pressurized water at 288°C, grain boundary rates of corrosion may be higher than in MgCl2 solutions at 154°C although this has not been shown. In addition, thicker films form at the higher temperature. Both factors could be expected to promote intergranular cracking and thereby make transitions to transgranular cracking less likely, as is found to be the case. Arguments have been advanced that all transgranular stress corrosion fractures have aspects similar to cleavage fracture (Bursle and Pugh, 1979). While it is not universally agreed that they all do derive from embrittlement by absorbed hydrogen or from any other embrittling reaction, a simple picture is emerging which at least needs to be examined. There is common agreement that practically all susceptible alloys exhibit an active path intergranular mechanism of cracking. What is not agreed is whether such a mechanism causes the cleavage-like fractures of transgranular cracking. In titanium and zirconium alloys the evidence is almost conclusive that hydrogen causes the transgranular fractures observed in environments capable of liberating hydrogen, and this appears likely for magnesium alloys also (Chakrapani and Pugh, 1976). In austenitic stainless steels strong

STRESS CORROSION CRACKING

161

evidence in support of hydrogen causing transgranular fracture has emerged recently, as described above. For brasses, evidence in missing. For mild steels, the rare evidence of the transgranular mode in stress corrosion failures does not allow a definite conclusion to be drawn. For aluminium alloys of high strength transitions to transgranular fracture are uncommon but this mode appears to be caused by hydrogen. In addition hydrogen is probably responsible for the major proportion (if not all) of the intergranular fracture. B. Crack tip chemistry The relatively simple description given in section III can be expected to be more complicated in reality but the general picture is probably close to being correct. Crack propagation occurs as the result of the interaction of a number of kinetic processes. Mechanically, the volume of metal behind the crack tip is undergoing a creep process, one effect of which is to produce fresh metal areas. Electrochemically, a repassivation process is operative, rendering the freshly exposed metal surface inert. The dissolved metal is undergoing hydrolysis and contributing to the composition of that volume of solution at the crack tip. Hydrogen ions are being discharged, with a proportion going into the metal. These various processes, some of which consist of a number of component reactions, e.g. hydrogen evolution and repassivation, are not likely to be acting independently of each other. As repassivation occurs, for example, the rate of hydrogen evolution will fall. As dissolution occurs the removal of dislocations may result locally in increased plastic flow, i.e. an enhanced creep rate. While it must be recognized that such complexities may define uniquely what really constitutes a stress corrosion situation, there is so little known about most of these areas of interaction and all that can be discussed is the simple model of creep and repassivation. The overall effect of these various events is to establish a balance at the crack tip characterized by a solution that has a narrow range of pH and composition, which together determine the repassivation rate. The pH values of solutions analysed from crack tips has already been described in section III. While in practice these values usually represent a fall in pH values from those of the bulk solutions in which the majority of service failures occur, such values are also encountered if the bulk solution pH is lower. For austenitic stainless steels in MgCl 2 solutions boiling at 125°C, for example, it has been shown that crack propagation occurs only when the pH of the corrodent within the crack lies in the pH range of 1 -2-2*5 as measured at room temperature

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(Baker et ai, 1970). A solution with an initial pH value lower than that range exhibited a rise in pH after exposure to stainless steel to within the specified range which was considered to be critical for stress corrosion because within this range a corrosion resistant protective film is formed. If a complexing agent such as glycerine or glycol, which prevents the formation of the protective film, was added to the solution, general corrosion occurred but no cracking, even if the pH value was within the critical range. Such a result emphasizes the critical role of the repassivation rate in controlling the crack propagation process as originally analysed by Scully (1967) and discussed above in relation to Fig. 6. C. The constant charge criterion The importance of the relationship between crack tip strain-rate and repassivation rate in controlling crack propagation has led to the Constant Charge Criterion (Scully, 1975). The charge passed during one increment of cracking in one grain, as drawn schematically in Fig. 7 has been hypothesized as being of a minimum value, Q*m. Propagation is considered conceptually to consist of a series of transients with the current occurring as a series of increments corresponding to increases in crack length. The value of the current is related to the crack tip strain-rate since the higher the strain-rate the more slowly will repassivation occur. This relationship is of considerable significance since the charge flowing is dependent both upon the creep rate and the repassivation rate. Raising the former by, for example, increasing the applied stress, and lowering the latter by, for example, raising the electrode potential (in many examples) will have the same effect of increasing the charge flowing in a given time. In Fig. 38 the relationship between strain-rate and current is drawn schematically for one transient under constant environmental conditions, e.g. potentiostatic control. The maximum current, /L, will be determined by electrochemical and hydrodynamic factors and may not be observed at low strain-rates. If the strain-rate falls to εΓ before Q*in has passed crack arrest will occur. From what has been indicated, this will happen, for a given solution, when the applied stress is unable to maintain the creep-rate above εΓ for sufficient time. For a given stress it will also happen if the repassivation rate is varied, for example, by inhibitor addition, or by lowering the potential. The occurrence of έτ is drawn in Fig. 39. In Fig. 40 crack propagation is drawn schematically as a series of transients under (a) constant stress conditions, and (b) constant load

Time

FIG. 38. Schematic diagram of one propagating transient. Initially there is a high strain-rate which diminishes rapidly. The strain-rate determines the value of the current flowing which also therefore diminishes. (Scully, 1975.)

Time

FIG. 39. Schematic diagram similar to Fig. 38. The strain rate falls to a value at which repassivation occurs, designed έΓ. This occurs before a minimum value of charge has passed, gmin, and crack arrest occurs.

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conditions, with each transient initiating the successive transient. g*in passes before έτ is reached. Under constant stress conditions each transient occurs within the same interval and the velocity, v, is constant: O*.

(4)

where / is the time interval between successive slip events and K is a constant. Under constant load conditions the stress will increase after

Time

FIG. 40. Schematic diagram of crack propagation during three increments of growth. Under constant stress conditions (a), the propagation rate is constant since there is no increase in overall strain-rate as the crack grows. This can be compared with constant load conditions (b), under which the propagation rate increases because the minimum charge passes in increasingly reduced time intervals because the overall strain-rate is increasing as the crack grows. Ultimately a maximum propagation rate is reached corresponding to an upper limiting current. (Scully, 1975.)

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each increment with a gradual increase in strain-rate which will cause ö*in to pass in increasingly shorter time intervals. This has been analysed as representing Region I cracking, with the velocity increasing exponentially as the stress increases (Scully, 1975). As the stress increases the maximum current, zL, will eventually occupy at the time so that Q*in = t x iL, at which point the velocity becomes constant and independent of increasing stress, corresponding to Region II cracking. The distance between successive film rupture events, Lc, should be that of step plane spacings which can be as high as 1 μπι. For mild steel in NaOH solutions it has been calculated as being c. 10 nm (Diegle and Vermilyea, 1975). Newman (1981) has pointed out that to reconcile such a large difference requires either that film rupture can occur between step events while the metal is straining elastically or that the intergranular crack path can act as a source of dislocations and produce a much finer slip pattern. Vermilyea (1972) had proposed such a model of the first kind, with film rupture occurring after repassivation was complete because of increasing elastic strain in the film. This is a different model from that drawn in Fig. 40 in which each transient initiates its successor before repassivation occurs. With such a wide range of cracking and repassivation phenomena there would appear to be a range of possibilities. Crack arrest or non-initiation can be observed in many systems subject to slow strain-rate conditions, however, corresponding to the crack tip strain-rate being ^έΓ> and manifesting as threshold stresses or stress intensities. Where very thin films exist these will be periodically fractured under such conditions, e.g. Ti alloys in aqueous NaCl (Scully and Adepoju, 1977). What is important is not the fracture but the repair in relation to the rate of fresh metal creation. Perhaps the Vermilyea model may find application in systems developing relatively thick films, for which it was developed. The concept of a Constant Charge has been developed also by Newman (1975) who has applied it to the stress corrosion of a turbine steel and a mild steel in NaOH. Both crack propagation and repassivation rates were determined as a function of electrode potential. The results were fully explained with reference to a film formation/film rupture mechanism, where crack growth is equivalent to the metal lost at the crack tip during film formation. The distance, Lc, was found to be a function of the physical properties of the steel and the applied stress intensity but was independent of the environment, whereas the propagation rate was dependent upon the environment. The experimental work included the assessment of the effectiveness of inhibitors on the cracking process by examining their effect

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upon repassivation. The model has been developed with the current being responsible only for film formation but it has been extended to allow for a proportion of the current being consumed in metal dissolution. In that situation Lc is not independent of the environment and tends to increase as the propagation rate increases. In the model of Scully (1975) no such distinction was made. The anodic current was considered to include both dissolution and film formation effects. From what has been described above, metal dissolution and hydrolysis is an essential feature of many stress corrosion systems, and the dissolution component was considered to be of major importance (Scully, 1975). The dynamical aspect of stress corrosion cracking processes has been further described in relation to the upper region of Fig. 8 and the similarities between pitting and cracking solutions in halide environments (Scully, 1980c). Whereas pitting occurs on a static surface, cracking occurs just below the pitting solution and it is argued that the solution composition is maintained in its pitting-like composition by the crack tip strain-rate which provides a continuous source of metal cations. A balance will be determined by the strain-rate and repassivation rate as is drawn schematically in Fig. 41 which attempts

Inhibitor

Increasing inhibitor

Increasing chloride Viscosity

STRESS CORROSION CRACKING

Halide Decreasing pH

Pitting potential

Increasing inhibitor Halide

Ductile failure Strain rate

FIG. 41. Schematic diagram of the relationship between electrode potential and crack tip strain rate, έ, for a solution in which passivation is possible. The boundaries are determined by solution composition, particularly inhibitors, halide ion concentration, pH and solution viscosity.

167

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to define the zone within which a solution able to cause cracking could be maintained. As the potential is lowered away from the pitting potential the repassivation rate increases and a higher strain-rate is required in order to maintain the solution composition. The upper boundary of strain-rate corresponds to conditions for maximum crack velocity which is potential-independent above a minimum value for some systems, e.g. Al (Speidel, 1971) and Ti alloys (Feeney and Blackburn, 1971). As the pH is lowered a constant velocity is observed and the boundaries of Fig. 41 will change to those drawn in Fig. 42. At the lower boundary of strain-rate cracking is replaced by corrosion rather than by repassivation. In the absence of έτ this transition will be a threshold for stress corrosion cracking. It may be difficult to recognize if the propagation rate is low and if intergranular cracking is replaced by intergranular corrosion. Such a transition would be shown up by experiments of the type summarized in Fig. 20.

Corrosion and stress corrosion cracking

Corrosion

Stress corrosion cracking

Ductile failure

Strain rate

FIG. 42. Schematic diagram of the relationship between electrode potential and crack tip strain-rate, έ, for a solution in which passivation cannot occur. At very low values of r, corresponding to tests lasting for a long period of time, corrosion occurs as on unstressed specimens, e.g. general corrosion or intergranular corrosion. This is also likely to be observed as the potential is raised in the noble direction. At active potentials cathodic protection may be achieved. Over a wide range of potentials the crack propagation rate will be constant. Because of the corrosiveness of the solution the boundary between corrosion and SCC may be difficult to discern. The upper boundary may have a slight positive slope.

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D. Fluctuating load conditions With increasing interest in Constant Extension Rate Testing, attention has been turned towards applying results to service conditions and to considering strain-rate effects in relation to practical operations. Whereas the time-to-failure tests summarized in Figs 1 and 2 employ constant load conditions, many engineering installations and structures are subjected to loads which vary during their use, e.g. vessels that are filled and emptied, or vehicles travelling in water. The origin of such considerations stems very much from the work of Parkins on the stress corrosion cracking of pipeline steels in HCO^/COJ solutions (Parkins, 1977a). Under static loading conditions a threshold stress was observed in solutions of pH 9-3 at temperatures of 75°C which was higher than the working pressure of the pipeline. Stress corrosion failure would not therefore be expected. If, however, the static load was varied slightly, and at a low frequency, failure occurred at much lower stresses, well below that caused by the internal gas pressure. This imposition of a frequency was entirely consistent with periodic increases in the gas pressure caused by passage through pumping stations. Crack propagation may occur only while the stress is increasing and only over that period of time within which the crack tip strain-rate has a suitable range of values able to sustain propagation, as has been discussed above. Such considerations can be of great importance, particularly in relation to service failures. When such a failure is being analysed, it must be borne in mind that the observed crack may have been growing for years, though intermittently, during the attainment of a crucial set of conditions which may not occur often. On a daily basis, crack propagation may have occurred during certain periods of emptying or filling procedures of vessels or pipelines. On an annual basis, propagation may occur under certain weather conditions or during an annual plant shutdown. The important feature is that the effect of a fluctuating stress is to increase the operative crack tip strain-rate. Evans and Parkins (1976), for example, have shown that load cycling can promote additional creep beyond that observed with static loading, with a strain increment occurring with each cycle. Propagation may occur, therefore, because the effective crack tip strain-rate has increased above the value ér referred to above. Such discussions have focused increased attention upon the use of Constant Extension Rate Tests since such experiments provide the possibility of reproducing under laboratory conditions the exact conditions that are likely to be encountered in service. Where stress or strain fluctuations

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are employed the possibility of inducing a corrosion fatigue effect must be considered. The transition between stress corrosion and corrosion fatigue has been examined by Parkins in pipeline steels (Parkins, 1977b). VI. Conclusion Stress corrosion cracking is an extensive subject incorporating electrochemistry, physical metallurgy and fracture mechanics and it occurs in many different alloy systems and environments. The principal body of experimental results constitutes a substantial publication (Staehle and Speidel, 1982). Any presentation is likely therefore to contain differences in emphasis. For example, no mention has been made in the above presentation of the mechanism of stress sorption (Uhlig, 1960) which attempts to explain cracking as arising from the fracture of stressed bonds as a result of the adsorption of a species from the solution at the crack tip. It is not altogether clear that it occurs during cracking in aqueous media. The stress corrosion of high strength steels has also not been described. There is much agreement that hydrogen embrittlement causes failure. The different affects of heat treatment and alloying are mainly concerned with metallurgical aspects of hydrogen species and dislocations, traps etc. (Thompson and Bernstein, 1980). Another omission is the modelling work of Krafft who has provided quantitative models of stress corrosion and corrosion fatigue derived from a formal model of tensile ligaments at the crack tip which are deforming and dissolving simultaneously (Krafft and Mulherin, 1969). The description in this chapter has been very much confined to phenomenological aspects of cracking rather than formal models. Readers may note other omissions in addition to the three indicated here. What has been attempted is a continuous modern picture of the propagation process in relation to experimental results as they are presently understood. Any account of the present understanding of stress corrosion cracking mechanisms is bound to be rather simplified since there are so many questions that cannot be answered at present because of a lack of data. Equally, there are probably relevant questions that have no yet been asked. The physical and chemical nature of the crack tip have not yet been determined or observed and it is in this region that the events of initial importance occur. The particular combination of morphological developments, adsorption of species at crack tip surfaces, hydrogen ion discharge and repassivation phenomena that

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cause stress corrosion crack propagation is unlikely to be easily unravelled, but at least some pointers exist and the picture described above is unlikely to be radically changed in the years to come. Instead, specific quantitative data will be produced and more realistic quantitative models will be developed. This is already occurring. The relationship between strain-rate and repassivation rate has been considered in detail (Scully, 1980). Predictive approaches are emerging from Constant Extension Rate Test experiments and slow potential scan-rates. There appear to be fewer areas of disagreement among workers than there were a few years ago. Stress corrosion cracking is a complicated phenomenon and, in addition, systems differ from each other; yet there is reason for hope and optimism that mechanisms are nearer to being clarified and problems anticipated.

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