Corrosion fatigue

Materials Science Engineering, 25 ( 1 9 7 6 ) 193 - 200

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© Elsevier S e q u o i a S.A., L a u s a n n e - - P r i n t e d in t h e N e t h e r l a n d s

Corrosion Fatigue

R. M. P E L L O U X , R. E. S T O L T Z a n d J. A. M O S K O V I T Z

Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139 (U.S.A.)

INTRODUCTION

Extensive research on corrosion fatigue over the past decade has produced a wealth of new information and theories but it has also raised a large number of unanswered questions. Although we certainly know much more now than we did ten years ago, we are also in a much better position to realize how much we do not know. For instance, we have clearly accepted the fact t h a t corrosion fatigue cannot be accounted for by the simplistic superposition of the processes of fatigue and stress corrosion cracking (SCC); however, we cannot ignore the similarities between the environmental effects in corrosion fatigue and in stress corrosion cracking. We have nicely covered our ignorance by describing corrosion fatigue as being due to a synergistic interaction between corrosion and fatigue. This interaction is k n o w n to take place b o t h during the initiation and the propagation stages of a fatigue crack. Although the exact boundaries between the initiation and the propagation stages are illdefined, we will deal with these two stages separately in order to review what we really do and do n o t know about corrosion fatigue.

CRACK INITIATION

The process of corrosion fatigue crack initiation, which has been reviewed extensively by Laird and Duquette [1], can be described most simply as the perturbation of the normal fatigue crack initiation mechanisms by environmental interactions, so that any unknowns pertaining to fatigue crack initiation must certainly apply to corrosion fatigue. For instance, we still do n o t know why, in an inert environment, fatigue failure

of a smooth bar can be indefinitely postponed below a certain cyclic stress amplitude. This lack of understanding makes it even more difficult to explain why there is no stress amplitude limit below which failure does not eventually occur in corrosion fatigue. The dislocation arrangements leading to the formation of a fatigue crack have been studied extensively [2 - 5], but mostly in the simple cases of pure metal single crystals. These studies, which are reviewed in detail by Grosskreutz and Mughrabi [6], show that one should differentiate between the dislocation structures in the bulk and in the persistent slip bands (PSB). In the bulk, the dislocations are arranged in veins or bands of dipoles, or in cell walls, depending upon the cyclic strain amplitude. This dislocation arrangement has been observed to exist in the bulk away from the PSB, but also within 500 - 1000 A of the free surface. Within the PSB, the dislocations are arranged in a welldefined ladder or cell structure. It is important to note that it has been reported [5] that the PSB are formed n o t only from newly generated dislocations but also by transfer of the matrix dislocations into the PBS structure. According to Basinski e t al. [7], this transfer can occur rapidly over a few cycles because the matrix dislocation structure is on the b o u n d a r y of stability. The demonstrated formation of PSB in surface layers which are initially free from dislocations appears to conflict with the findings of Kramer and Kumar [8] who have claimed the existence of a surface debris layer. The important point, however, is that dislocation arrangements associated with PSB, which eventually form intrusions and extrusions and act as crack initiation sites, and the mechanisms responsible for this process, have been characterized only for the simplest, oxidefree single crystals.

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Studies by Greenfield e t al. [9 - 11] have shown that surface deposition of a thin layer (500 A) of metal followed by surface alloying can greatly change the near-surface plastic deformation characteristics of pure copper by inducing the formation of a network of accommodation dislocations which act as barriers to slip activity. The resulting dislocation tangle beneath the surface layer disperses the slip and prevents localization, thereby delaying crack initiation considerably Whether surface oxide layers, which exist on many metals and alloys, can act as such barriers is unclear, as is the very nature of many of these oxide layers. Environmental interactions are, b y nature, surface interactions, so it is the aforementioned near-surface dislocation phenomena which are strongly affected by the environment. An active environment can either affect the formation of PSB, or affect the mechanisms by which a crack develops from these PSB. Much of the uncertainty a b o u t corrosion fatigue lies on this point. Howard and Pyle [12] have shown a relationship between PSB spacings and electrochemical dissolution rates of a metal surface, with PSB spacing decreasing with a decreasing dissolution rate. If one considers a surface oxide layer as a barrier to slip, similar to the coatings applied b y Greenfield e t al. [9 - 11], it follows that in the presence of an oxide layer, a case for which the dissolution rate will be low, PSB spacing will be small because slip localization is inhibited. By raising the dissolution rate, or by making the oxide film thinner, the barriers to slip are reduced, slip localization is favored, and PSB spacing increases greatly. Studies on copper single crystals have shown that all of the imposed plastic strain is taken up by the PSB regions [13]. Thus, if the PSB spacing is increased by environmental effects, each PSB must carry a larger local plastic strain, leading to accelerated crack initiation. Additionally, the environment can alter the stability of the near-surface matrix dislocation structures and cause the matrix dislocations to be transferred more easily into the PSB, further accelerating PSB formation. Alternatively, it can be said that PSB form independently of any environmental effects, and that the role of the environment is to accelerate the process by which PSB lead to fatigue crack initiation through intrusion-

'(a)

(b) Fig. 1. Effect of surface dissolution on PSB development. A o = 351.6 MPa (51 ksi). (a) In air, PSB development is insufficient to promote crystallographic cracking; (b) in acidic chloride solution, greater PSB development produces strongly crystallographic cracking.

extrusion formation. This can occur b y three processes: (1) preferential dissolution at PSB because the PSB is anodic to the matrix; (2) film-rupture by emerging PSB causing the exposure of fresh metal; (3) localized dissolution resulting in marked local changes in dislocation arrangement, affecting the intrusion-extrusion development within the PSB. Pyle, Rollins and Howard [14, 15] have proposed, on the basis of studies of dissolution transients associated with cyclic plastic straining, an initiation mechanism of slipdissolution-reverse slip at PSB to explain the relationship between PSB spacing and

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dissolution rate. Moskovitz [16], however, has shown t h a t at a constant alternating stress, PSB in austenitic-ferritic stainless steels develop more rapidly in the presence of a corrosive environment (Fig. 1). This observation is consistent with the process of dissolution of the surface oxide layer resulting in enhanced PSB formation. The effect of surface dissolution on near-surface dislocation structures has recently been studied by Duquette e t al. [17, 18] in an a t t e m p t to show t h a t changes in near-surface dislocation arrangements are caused by environmental interactions. Duquette found in OFHC copper that there was a difference in the near-surface dislocation structure between specimens tested in air and in NaC1 solution. He also inferred t h a t this difference was large enough to have an effect on PSB formation. However, up to now, no study has been able to show definitively which specific environmental interaction controls the PSBintrusion-extrusion processes. Much of the work on environmental interactions mentioned above was influenced by the results of studies by Lee and Uhlig [19], Duquette and Uhlig [20], Revie [21], and Asphahani [22]. All of these studies have shown the existence of a critical anodic current level, below which there are no effects of electrochemical dissolution on fatigue life. This critical current has been interpreted as the a m o u n t of current necessary to dissolve away dislocation pile-ups at slip steps where intrusion-extrusion formation was impeded [17, 20]. Another interpretation of the current is t h a t it produces sufficient dissolution on freshly exposed slip step faces so t h a t upon slip reversal a n o t c h of critical size is formed [14, 15]. It also has been suggested t h a t preferential dissolution of active slip bands [17] leads to intensification of slip, or t h a t the dissolution is associated with the production of divacancies which induce climb in sessile dislocations [21]. This concept of a critical current is the key to the understanding of corrosion fatigue, and any proposed theory of crack initiation must be able to account for its existence. The main problem involved in resolving this question is to determine w h e r e the dissolution takes place, and to c o n v e r t a critical c u r r e n t value to a critical local c u r r e n t d e n s i t y , so t h a t the

dissolution rate can be given a physical significance. The gap between our understanding of bulk corrosion phenomenon, characterized by an overall current density, and the microprocesses occurring at selected surface sites, remains the largest stumbling block in modeling corrosion fatigue crack initiation. Further complications arise in the understanding of the initiation process when we consider cases other than the simple electrochemical dissolution of a smooth surface. Prowse and Wayman [23] propose t h a t hydrogen embrittlement effects are responsible for accelerated corrosion fatigue failure in a medium carbon steel. Effects such as this make it difficult to differentiate corrosion fatigue from other environmentally assisted fracture modes, such as stress corrosion cracking. The superposition of a mean stress on cyclic stresses produces additional complications, as SCC effects become much more significant in the presence of a mean tensile stress. Consideration should also be given to the role of the near-surface state of stress (ratio of tensile to shear stress) on the corrosion fatigue mechanisms. This can be done by studies of notched bars, such as those performed by Barsom and McNicol [24], in which the number of cycles to initiate a fatigue crack was found to be a function of the notch radius and notch depth. Similar studies have yet to be done for corrosion fatigue. Finally, corrosion effects can change the relative fraction of life spent in initiation and in propagation of a fatigue crack. As in the case of the notched bar tests [24], it becomes increasingly difficult to differentiate between crack initiation and very low growth rate crack propagation. Further experiments on the growth of very short cracks, and in the threshold region of fatigue crack growth are needed in order to separate the influence of corrosion on the transition from crack initiation to crack propagation.

FATIGUE CRACK PROPAGATION

We shall now consider our understanding of corrosion fatigue crack propagation (CFCP), where a growing fatigue crack is subjected to cyclic loading in an active environment. There are a number of

196 similarities and differences between the mechanisms of crack propagation and those of crack initiation. As in the case of initiation, we still must consider the processes of oxide or film growth, passivation, and diffusion, and their overall effect on the near-surface dislocation arrangement and local mechanical properties. By contrast to crack initiation, however, the site of the chemical-mechanical interaction is restricted to the crack-tip region where the local strains are much higher than those in the PSB. During cyclic stressing, fresh metal surfaces are constantly being exposed at the crack tip, so that transient oxidation or passivation effects of the order of milliseconds become very important. The amount of surface areas exposed are related to the crack tip opening displacements which can be as small as a few hundred angstroms. The local chemistry near the crack tip is quite different from that in the bulk of the corrosive solution so that the interpretation of externally measured corrosion parameters is difficult. Lastly, we are concerned with the acceleration of crack growth rates by the corrosive environment over and above that observed in an "inert" environment. Thus, the best we can hope for in a practical sense is to reduce the crack growth rate rather than indefinitely to postpone fatigue failure, as is possible in the initiation stage. The simplest approach to CFCP is to consider that crack growth is due to the superposition of pure fatigue and SCC, ignoring any synergistic effects. Given a threshold stress intensity for SCC, K~scc, one can calculate the time spent above K, scc during each stress cycle and add the component of corrosion crack growth to the mechanical fatigue c o m p o n e n t [25, 26]. This assumes that Kiscc is unaffected by the loading rate, which is n o t always true [27]. Alternatively, it has been suggested that only when the crack growth rates due to fatigue and SCC. are equal can one apply a superposition model [28]. This often occurs at stress intensities above K~scc. It is obvious that these models can be used only for materials which are stress corrosion susceptible. While the superposition approach has had some success in describing empirically the crack growth rate in selected alloys, it does n o t help us in the understanding of the basic mechanisms.

Environmentally accelerated fatigue crack growth also occurs in materials and at stress intensities where no SCC is observed. It is in this regime, called "true corrosion fatigue", where the greatest number of questions still remain. Classical experiments with different stress waveforms [29 - 31] have shown that the major part of crack growth occurs during the rising portion of the stress wave, rather than at the peak stress. This would indicate that the key to the mechanisms of true corrosion fatigue lies in a dynamic interaction between the environment and the exposed metal surface undergoing plastic strain. In trying to assess our understanding of corrosion fatigue crack propagation we have divided the various proposed mechanisms into three general classes: (1} theories involving dissolution of metal at the tip, (2) theories involving essentially mechanical crack-tip effects, {3) theories involving changes in the local deformation character of the material at the crack tip. The theory of active dissolution, proposed by Ford and Hoar [32], states that in A1-Mg alloys, the increase in CFCP in chloride and sulfate solutions can be entirely accounted for b y the amount of material dissolved at the crack tip. The high local current transients necessary to account for dissolution rates have been measured by means of a scratching electrode test. The theory suggests that enhanced anodic reactions always increase the growth rate which, as we shall show, is not always true. The second class of theories suggests that the environment changes the nature of the oxide layer from ductile to brittle [33]. During straining, microcracks form in the oxide, leading to an increased local stress concentration and a mechanical increase in crack growth rate. A thick enough oxide layer can, in some cases, be beneficial, however, either b y blunting o u t the crack and preventing complete resharpening [34], or by wedging the crack open and reducing the effective stress intensity factor. Finally, the environment can affect the profile of the macroscopic crack front across the specimen, reducing the overall growth rate by crack branching.

197 The third category of mechanisms is based on environmentally induced changes in deformation character. Single crystal experiments by Latanision and co-workers [ 35, 36] have shown t h a t corrosion effects at the surface can change markedly the flow behavior of the entire crystal. Among the mechanisms that could be operative at the crack tip is t h a t suggested by Tyson and Alfred [37 ], where an adsorbed species reduces the cohesive strength of the material sufficiently to promote a cleavage mode of fracture. While such a mechanism may explain liquid metal embrittlement [38, 39], it is unlikely to apply to f.c.c, metals in other liquid environments because the very low ratio of critical shear to critical cleavage stress should always result in deformation by slip rather than cleavage. In another proposed theory the distribution o f slip at the crack tip can be altered by diffusion of some species from the surface into the plastic zone. Frandsen et al. [40] report a change from wavy or cellular to planar slip mode in a Ni-Cu alloy in the presence of hydrogen gas. This change of slip mode shifts the fracture path from transgranular to intergranular, with a c o n c o m i t a n t increase in CFCP. By contrast, in aqueous solutions or at high temperatures, a coherent oxide may prevent the emergence of sharp slip steps at the fresh metal surface of a planar slip material, thus diffusing the slip and reducing the crack growth rate. Such a mechanism has been proposed by Gell and Duquette [34] for the fatigue behavior of single crystals of nickel-base superalloys. Finally, the environment may affect the kinetics of repassivation, which, in turn, affects the flow properties of the crack tip. Anomalous frequency effects in Ti alloys (decrease in growth rates with a decrease in frequency) have been explained by noting that, at lower frequencies, more time is allowed for repassivation at the crack tip, the authors concluding t h a t repassivation either minimizes the dissolution rate or leads to slip dispersal [27]. In summary, it is difficult, if n o t impossible, to isolate which of the above mechanisms controls the corrosion fatigue behavior of an alloy in a given environment. The following example will serve to illustrate how much progress has been made in understanding

Fig. 2. Fatigue fracture surface of 7075-T6 alloy tested at 1 Hz in 3.5% NaCI solution. Brittle striations (A) form during cathodic polarization. Ductile striations (B) form during anodic polarization. (Imposed current was 0.5 mA: 10 s cathodic, 30 s anodic.)

Fig. 3. Details of brittle and ductile striations in 7075-T6. Brittle striations are generally flat with "river markings" perpendicular to striation front. Ductile striations are smooth ripples. The brittle striations are ~ 5 times wider than the ductile striations. CFCP and where the difficulties still lie. A1-Zn-Mg (7075) alloys are susceptible to corrosion fatigue in chloride solutions. Not only does the presence of C1- ions accelerate the growth rate ten-fold over t h a t in air, it also changes the type of fatigue striations formed [41]. In air, the striations are evenlyspaced ripples on the fracture surface. They are termed " d u c t i l e " striations because their overall appearance suggests a large a m o u n t of plastic blunting at the crack tip. On the other

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hand, "brittle" striations, which form in chloride solutions, are much more uneven in their appearance and have cleavage-like features. Figures 2 and 3 are from a sample of 7075-T6, fatigue tested in a 3.5% NaCl solution at a frequency of 1 Hz. Brittle striations would normally form under these conditions: however, we have imposed on the crack surfaces a direct current of 0.5 mA alternately polarizing the sample anodically and cathodically against a Pt electrode. Specifically, the cathodic current is applied for 10 s (10 stress cycles) while the anodic current is applied for 30 s C30 stress cycles). During anodic polarization the fatigue process changes from normal "brittle" striations (at A in Fig. 2) to ductile striations such as would be present if testing were done in air Cat B in Fig. 2). Cathodic polarization does not change the type of features normally present in chloride solutions. The change in mode is instantaneous (less than 1 s) with change in polarity (there are 10 brittle and 30 ductile striations). It is also interesting that Ca) the sharp transition occurs only in a selected, well oriented grain, (b) the ductile-brittle transition is very smooth, while the brittle-ductile transition is highly irregular (Fig. 3). We shall now consider which of the theories can explain this behavior. The dissolution model is not applicable because the local growth rate is reduced by anodic polarization. The mechanical effect of the oxide layer is a more plausible explanation. Under cathodic conditions, local chlorine enhanced pitting occurs at sites along the crack front, causing the crack to advance unevenly at each cycle. The oxide does n o t appear to be sufficiently thick to protect against local pitting attack. Under anodic conditions, the oxide is thick enough to prevent pitting, thus crack advance is smooth along the front and a great deal of crack tip blunting occurs. The mechanical behavior of the oxide film alone, however, cannot explain the change in fracture plane orientation. A cleavage mode of fracture, originally proposed for "brittle" striations, has been discounted because similar alloys exhibit crystallographic fatigue features, even in the absence of a corrosive environment [42]. It is more likely that during anodic polarization the oxide disperses the sharply defined slip bands that form ahead of the crack tip. Under cathodic

conditions, these sharp flow bands lead to a strongly crystallographic fracture path, as is predicted in various models of fatigue crack growth [43, 44]. Thus, a combination of mechanical oxide effects and a change in slip character best describes the corrosion fatigue crack propagation behavior of A1-Zn-Mg alloys. We still do n o t understand completely, however, the specific chemical reactions that take place at the crack tip leading to the ductile-brittle striation transition. Further studies are needed on the various oxides that form in A1-Zn-Mg alloys and on the mechanical strength of these oxides. Such studies would also be valuable in other alloy systems in order to differentiate between the mechanical and substructure effects leading to corrosion fatigue.

DISCUSSION

Following this review of what we do and do n o t k n o w about corrosion fatigue, it is fair to ask what should we do next? There is no d o u b t that more work on the micromechanisms of plastic deformation and fracture is badly needed. This work will be difficult b u t challenging, because it will require the use of advanced tools and techniques to differentiate between the mechanisms of dry fatigue and corrosion fatigue. Some of the more worthwhile research topics are mentioned below. C1) Studies of the near-surface dislocation arrangement as a function of surface films and active environments. The use of high voltage electron microscopy combined with the pinning of dislocation arrangements by a neutron flux seems to be the best w a y to face this difficult problem. (2) Analyses of corrosion fatigue fracture surfaces by Scanning Auger Spectroscopy. This technique should help us to characterize the ions which are adsorbed at the crack tip during the corrosion-fracture process. (3) Measurements of corrosion current transients on freshly broken wires, or scratching electrode tests with different solution chemistries. The data would permit a quantitative evaluation of the dissolution rates in the vicinity of emerging slip steps, within persistent slip bands, or at the crack tip.

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(4) Fractographic studies of the transitions in fracture modes from dry fatigue to corrosion fatigue crack propagation (and vice-versa). The detailed fracture features will provide a better description of the micromechanics parameters controlling fracture at the crack tip. In summary, advances in corrosion fatigue research will require an interdisciplinary approach between experts on microplasticity and dislocation theory, corrosion and electrochemistry, and fracture mechanics. A better understanding of the problem is not only of academic interest, but it is urgently needed by the design engineers, who are trying to minimize corrosion fatigue and to improve the reliability of structural components operating for lives as long as thirty years in more and more corrosive environments.

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