Corrosion fatigue mechanisms in b.c.c. stainless steels

Corrosion fatigue mechanisms in b.c.c. stainless steels

Acra metal/. Vol. 35, No. 8, pp. 2105-2113. Printed in Great Britain. All rights reserved 1987 CORROSION Copyright FATIGUE MECHANISMS STAINLESS ST...

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Acra metal/. Vol. 35, No. 8, pp. 2105-2113. Printed in Great Britain. All rights reserved

1987

CORROSION

Copyright

FATIGUE MECHANISMS STAINLESS STEELS T. MAGNIN

OOOI-6160/87 $3.00 +O.OO ((‘8 1987 Pergamon Journals Ltd

IN B.C.C.

and L. COUDREUSE

Departement Materiaux, Ecole des Mines. 158. tours Fauriel, 42023 Saint-Etienne Cedex, France (Received 10 July 1986; in revised form 18 November 1986) Abstract-The

influence of a 3.5% NaCl solution on the cyclic plastic deformation of a b.c.c. Fe-26CrrlMo alloy is analysed as a function of the applied electrochemical potential E, taking into account the dislocation behaviour, the formation of microcracks and the evolution of the cyclic corrosion current transients. Three different kinds of damage mechanisms have been pointed out: (i) in the cathodic region (E < - 300 mV/SCE), the localization of the hydrogen reduction is favoured when microcracks are mechanically formed. This induces a very marked decrease of the fatigue life at high strain amplitudes, but this deleterious effect is reduced at low strain amplitudes for which the microcrack formation is delayed; (ii) in the passive region for - 300 < E < + 500 mV/SCE, the damage mechanisms are related to the localization of the anodic dissolution due to a depassivationrepassivation process. A critical strain rate range (- 10-j s-‘) for which this localization and the corresponding acceleration of the microcrack formation are maximum is encountered; (iii) in the passive region for E > + 500 mV/SCE, the cyclic strain enhances the formation of pits which induce an early formation of microcracks. The study of the microcracking process and its evolution is a key to specify the physical mechanisms by which an aqueous corrosive solution can affect the fatigue life of b.c.c. stainless steels according to the applied electrochemical potential. R&sum&-L’influence dune solution a 35% de NaCl sur le comportement en deformation plastique cyclique dun alliage C.C. Fee26CrrlMo est analysie en fonction du potentiel Clectrochimique applique. Cette etude prend en compte a la fois le comportement des dislocations, la formation des microfissures ainsi que i’evolution des transitoires de courant. Trois types differents de mecanismes sont soulignes: (i) dans le domaine cathodique (E < - 300 mV/ECS), la localisation de la reduction de l’hydrogene est favorisee quand des microfissures dues a la fatigue seule se forment. Ceci conduit a une diminution importante de la durbe de vie aux fortes amplitudes de deformation; mais cet effet endommangeant est reduit aux faibles amplitudes de deformation pour lesquelles la formation de microfissures intervient tardivement; (ii) dans le domaine passif pour - 300 < E < + 500 mV/ECS, l’endommagement est dfi a la localisation de la dissolution anodique like a un mecanisme de depassivation-repassivation cyclique. Un domaine critique de vitesse de deformation (1 10e3 s-‘) pour lequel cette localisation et l’accileration de la microfissuration qui lui est lite sont maximum est mis en evidence; (iii) dans le domaine passif pour E > 500 mV/SCE, la deformation cyclique induit la formation de piqures qui sont a l’origine de l’amorqage precoce de microfissures. L’etude de la microfissuration en surface constitue une cli pour comprendre les mecanismes physiques par lesquels une solution aqueuse peut reduire la durte de vie d’alliages inoxydables C.C., en fonction du potentiel ilectrochimique impose. Zusammenfassung-Der

Einflu5 einer 3,5%igen NaCl-Liisung auf die zyklische plastische Verformung der krz. Legierung Fe-26Cr-1 MO wurde in Abhangigkeit von dem angelegten elektrochemischen Potential E unter Berticksichtigung des Versetzungsverhaltens, der Bildung von Mikrorissen und der Entwicklung der zyklischen Transienten im Korrosionsstrom analysiert. Drei unterschiedliche Schadigungsmechanismen wurden gefunden: (i) Im katodischen Bereich (E < -300 mV/SCE) ist die Lokalisierung der Wasserstoffreduktion begiinstigt, wenn Mikrorisse mechanisch gebildet werden. Dadurch wird eine sehr deutliche Verktirzung der Ermiidungsstandzeit bei hohen Dehnungsamplituden verursacht. Dieser nachteilige Effekt ist allerdings bei niedrigen Dehnungsamplituden, bei denen die Mikrorigbildung verziigert ist, geringer. (ii) Im passiven Bereich - 300 < E < + 500 mV/SCE hangen die Schldigungsmechanismen mit der Lokalisierung der anodischen Aufliisung durch einen Entpassivierungs-Repassivierungsprozeg zusammen. Es findet sich ein kritischer Bereich der Dehnungsrate ( = 10-j SK’), bei dem diese Lokalisierung und die entsprechende Beschleunigung der MikroriBbildung am grogten ist. (iii) Im passiven Bereich E z + 500 mV/SCE verstarkt die zyklische Dehnung die Bildung von Griibchen, welche zu einer friihen Bildung von Mikrorissen fiihren. Die Untersuchung des Bildungsprozesses von Mikrorissen ist der Schliissel zum Auffinden der phsikalischen Mechanismen, mit denen eine wassrige korrosive Liisung die Ermtidungsstandzeit von krz. rostfreien Stahlen je nach dem angelegten elektrochemischen Potential

beeintrichtigen

kann.

1. INTRODUCTION

The reduction in fatigue resistance alloys in the presence of an aqueous AM

3S!L-L

of metals and corrosive solu-

tion is generally associated with a local interaction between a mechanical and an electrochemical damage. During corrosion-fatigue, the localization of the plastic strain due to fatigue can enhance local electro2105

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transients and the analysis of the surface roughness (particularly the microcrack formation). 2. EXPERIMENTAL

Fig. 1, Current transient characteristics at imposed potential (J = I/1.5 cm2, where I is the anodic current).

chemical damage which may be described in terms of dissolution and/or hydrogen reduction reactions. Moreover, the corrosion damage can likewise favor the localization of the cyclic strain [1,2]. Thus, the corrosion fatigue behaviour of pure metals and alloys in a given solution is closely dependent on both the plastic deformation mode which controls the strain localization during cycling [3], and the electrochemical potential which governs the kinetics of the corrosion reactions [4,5]. To study the corrosionfatigue mechanisms it is then necessary to perform tests for which the basic mechanical, metallurgical and electrochemical parameters can be imposed or controlled. Thus, even if many papers have shown the effect of a corrosive solution on the fatigue properties of metals (for a review, see [li, A), only a few studies have focused on the mechanisms. Some performant electrochemical tests have been used to study the enhancement of the dissolution behaviour by the cyclic strain on mild steel and austenitic stainless steels in different solutions [4,5,8,9] but without any analysis of the dislocation behaviour and the plastic deformation mode. Inversely, some very nice studies have described the effect of a corrosive solution on the strain localization during the cyclic plastic deformation of pure copper [2, IO] taking into account the dislocation behaviour but only in the case of an anodic dissolution without any passive film. The main purpose of this paper is to analyse the b.c.c. mechanisms of a corrosion-fatigue Fe-26Cr-1Mo stainless steel in a 3.5% NaCl solution using low cycle fatigue tests at imposed plastic strain, taking into account both the dislocation behaviour and the evolution of the corrosion current transients for different imposed electrochemical potentials. So the corrosion fatigue behaviour of the ferritic stainless steel is examined in the domains of (i) passivity, (ii) pitting, (iii) hydrogen reduction. A particular attention is paid to analyse the evolution of the surface damage of the specimens using a correlation between the observation during cycling of the current

Corrosion-fatigue tests were carried out on a high purity ferritic stainless steel containing 25.526 wt% Cr, 0.94-l. 15 wt% MO and about 20 ppmC and 30 ppmN, in the form of single and poly-cristals. The heat treatment consists of a 1 h solution anneal at 1050°C and a waterquench. This alloy has been previously studied in cyclic deformation in air and the evolution of the dislocation microstructures has been determined [l 1, 121. A very sensitive testing equipment is used to simultaneously record the cyclic evolution of the mechanical and electrochemical parameters. This method has been described in detail elsewhere [13]. Symmetrical tension-compression tests were performed on a servo-hydraulic machine on smooth specimens under plastic strain control (+ 10e4 < fAf,/2 < 5 10e2), at constant strain rate i ( low5 s-’ < i < lo-* s-‘) in dry air and in an aerated and replenished 3.5% NaCl solution at pH 6. Some tests were also conducted in the desaerated solution, particularly when the applied potential corresponds to the cathodic region. The specimens, with a 5 mm diameter and a 10 mm gauge length, were partially coated with lacquer and Teflon a part from a window of 1.5 cm’ exposed to the corrosive solution. A sensitive potentiostatic system is used to impose the electrochemical potential E, referred to a saturated calomel electrode (SCE). During cycling at imposed potential, the current transients are recorded as a function of the applied strain and the number of cycles. Figure 1 schematises the evolution of the current density J (J = I/l .5 cm*, when Z is the recorded current) during one cycle. To quantify the perturbation of the metal-solution interface during cycling the parameters JT, J,, Jb and the quantity of electricity generated per cycle AQ are used and their evolution during cycling is followed as a function of the imposed mechanical and electrochemical conditions. Finally, to determine the surface roughness and to quantify the corrosion-fatigue damage, cellulose acetate replicas of the specimen surface are taken for different fractions (1~100%) of the fatigue life Ni corresponding both to the initiation of a 1 mm long macrocrack and to a rapid 1% decrease of the saturation stress. The replicas are then analyzed using optical and scaning electron microscopy. Histograms are drawn to present the number of cracks per unit area (1 mm*) and their length in pm for the different fractions of Ni. 3. RESULTS 3.1. Electrochemical behaviour without q&g The polarization curves for the Fe26Cr-lMo alloy in the aerated and desaerated 3.5% NaCl solution at pH 6 are given by the Fig. 2. The free

MAGNIN and ~OUDREUSE:

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CORROSION FATIGUE IN B.C.C. STAINLESS STEELS

L

-1000

Fig. 2. Polarization curves of the Fe-26Cr-1Mo alloy in the aerated and desaerated 3.5% NaCl solutions (293 K, dE/dt = 90 mV/h).

of the more mobile edge dislocations, and the above grain boundary effects are not encountered 1141. At low strain rates [Fig. 3(b), and for i < 10-4s-‘] the dislocation arrangement is quite different. The total dislocation density is higher, dislocation bundles composed of edge multipoles are well defined and oriented perpendicularly to the Burgers vector. The mobile dislocations between the bundles are curved and of mixed character, the mobilities of screw and edge dislocations being quite similar. The asymmetrical behaviour described for high strain rates is suppressed. Cross slip is favoured which promotes the formation of the well known persistent slip bands (PSB). Transgranuiar crack initiation is then observed in polycristals 1111. 3.3. Infltlence of the applied potential on the fatigue life in the aerated 3.5% NaCl solution

potentials are about & = - 150mV/SCE and - 300 mV/SCE for the aerated and desaerated solution respectively. The pitting potential is about +800 mV/SCE in each solution. The corrosionfatigue tests will be then performed (I) in the passive E= -3OOmV/S~E,~=~~,~= region (at + 200 mV/SCE, E = + 500 mV/SCE), studying the influence of the thickness and stability of the passive layer, (2) in the pitting region (+600 mV < E c 8SOmV/SCE), (3) in the cathodic region (at E = - 500 mV/SCE and E = - 800 mV/SCE). 3.2. Dislocation behaviour during cycling in dry air The cyclic plastic deformation process of the b.c.c. Fe-26Cr-1Mo alloy cycled at 300 K has been determined in previous studies [ll, 121. It is strongly dependent on the strain rate because the temperature To above which an athermai behaviour is approached (corresponding to a similar mobility of screw and edge dislocations 1141) lies in the vicinity of room temperature for the ferritic stainless steel. The Fig. 3 which indicates the dislocation substructures observed on the (211) slip plane of [Ol I] single crystais cycled at Ac,/2 = 4.10m3 during 10 cycles, clearly shows the influence of the strain rate li. At high strain rate [Fig. 3(a), and for i 3 1O-3 s-l], the dislocation arrangement consists of a substantial density of long straight screw dislocations and of small dislocation clusters which are distributed without any well defined order. The screw dislocations, the mobility of which is much lower than the mobility of edge dislocations, control the defo~ation (at Ac,/2 2 lo-‘). Because of the special core structure of screw dislocations in the b.c.c. lattice, dislocation glide often occurs on different slip systems between tension and compression [12, IS]. Pencil glide is induced and associated with a significant plastic strain irreversibility between tension and compression which promotes a shape change of the grains [14] and a strain localization at grain boundaries for polycristals [I I]. When AQ,/~ < 10s3, the strain is principally accommodated by the quasi reversible motion

Figure 4 clearly shows a marked effect (at Ac,/2 = 4.10e3 and t: = 10Y3s-‘) of the applied electrochemical potential E on the number of cycles N, corresponding to a rapid 1% decrease of the saturation stress u,. It can be observed on Fig. 4 that: (i) a reduction by a factor 2 of the fatigue life in air occurs at E = E0 (passive region); (ii) in the passive region, the higher is the potential, the longer the fatigue life until E = +500 mV/SCE. But at E = t600 mV/SCE, a new decrease of N, is observed; (iii) finally, a very marked reduction of N, occurs in the cathodic region (reduction by a factor 5 at E = -800 mV/SCE).

b

J’

Fig. 3. Dislocation substructures on the (%I) slip plane of [OlI] crystals cycled at &/2 = 4,10e3 for !O cycles at (a)

l

=2~1O-“s-’

and (b) g =2,10-‘~-~.

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FATIGUE

IN B.C.C. STAINLESS STEELS

10-z-

5 +\, Air

7

_; -

/

:

L

- 500mV

I

!/’

#lo-‘-





\

I4 Eo

26Cr-IMO 3.5 % NaCl (aerated)

J

ASP,4.lQ3,i.liy3<'

2

I -800

I

I

-500

0 E

(mV/SCE

I

500

1

I

IO'

I

IO'

IO4

I

IO'

Ni

Fig. 4. Influence of the applied electrochemical potential E on the fatigue lifetime N, in the aerated solution at At,/2 = 4.10-’ and 5 = IO-'s-‘.

Fig. 6. Influence of the electrochemical potential on the Coffin-Manson curves at i = 10m2s-i (aerated solution).

The corresponding crack initiation sites has been determined. For cathodic potentials and in the passive region at 300 mV < E < 500mV/SCE there is no influence of E: crack initiation is intergranular [Fig. 5(a)] as in air [l 11, which has been related to the shape change of the grains caused by the asymmetrical behaviour of the screw dislocations between tension and compression. In contrast, at E > 500mV/SCE the reduction of N, is associated with pitting (essentially at grain boundaries) as it can be Seen on Fig. 5(b). The effect of the imposed plastic strain amplitude on the fatigue life reduction is also dependent on the applied potential as it can be seen on the Fig. 6. In the passive region (at E = E,)the Coffin-Manson curve exhibits the same slope ( N -0.9), in the 3.5% NaCl solution and in air, even if crack initiation is intergranular at Ac,/2 > 1O-3 and transgranular at AcP/2 < 10 A3as shown and explained elsewhere [l 11. In contrast, the influence of At,/2 is sensitive for applied cathodic potentials: the scatter between the

fatigue life at E = - 500 mV and E = E,, decreases when AeJ2 decreases. Finally, the effect of strain rate is also dependent on the applied potential. It has been studied in detail at E = E,in a previous study [13]. In this case (Fig. 7) a critical strain range in the vicinity of 10m3s-’ is observed for which lifetimes are significantly reduced. Moreover, if at high strain rates (i N IO-*s-r) the effect of the corrosive solution is also noticed, no corrosion-fatigue occurs at i < 10e4 s-‘. Such a behaviour has been observed for -300 mV < E < + 500 mV. In contrast, whatever i, a reduction of the fatigue lifetimes is encountered for cathodic potentials (Fig. 7): at Ac,/2 = 4+10d3, Ni N 300-400 cycles for lo-’ SK’< i < lo-* SK’. 3.4. Influence of the applied potential on the fatigue damage evolution The analysis of the replicas of the specimen surface during cycling in air and in the corrosive solution at imposed potential leads to classify the cracks in four

Fig. 5. Crack initiation during corrosion fatigue at [email protected] = 4.10m3 and i = lo-’ SK’: (a) at E = 170,(b) at E = +600 mV/SCE (ptttmg is enhanced).

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i $

30-

10woo

1300

N Fig. 9. Influence of an applied cathodic potential (E = -500 mV) on the microcrack formation (AcJ2 = 4.10-‘, I = 10-j s-‘, aerated solution).

Fig. 7. Influence of strain rate on the fatigue lifetime N, at Acp/2= 4.10W3(aerated solution).

types as a function of their length I 1161: (if type I cracks for 1 < 50 pm (typically the grain boundary size); (ii) type II cracks for 50pm < 1 < 150 pm. These cracks form by micropropagation in surface of type I cracks 1161; (iii) type III cracks (for 150 pm < 1 < 500 hrn) which form by coalescence and micropropagation in surface of type I and type II cracks. They are generally oriented ~r~ndicularly to the stress axis; (iv) type IV cracks (for 500 pm < 1 < I mm) which lead to a macropropagation in volume. 1 or 2 type IV cracks are generally observed in our specimens. On Fig. 8, the number of cracks per mm2 of types I, II and III is indicated, as a function of the number of cycles N, in air and in the corrosive solution at E = E,, (passive region) for At,/2 = 4*10s3 and i = 10F3 s-l, It can be observed that: (i) in the corrosive solution the type I cracks initiate for the same number of cycles (N = 100) than in air, but for this value of N their number is much larger in the 3.5% NaCl solution; (ii) the type II and type III cracks can form more rapidly in the corrosive solution, which explains the reduction of the fatigue at E = E,. It can be also noticed that at N = Ni, the total number of type I, II and III cracks is identical in air and in the corrosive solution

Fig. 8. Evolution of the number of type I, II and III cracks during cycling at Ac,/2 = 4.10-l, 1 = IO--’ SC’ in air and at E = E0 (aerated solution).

(it must be emphasized that many cracks do not propagate). Nevertheless the damage leading to rupture is obtained at N = 700 cycles in the 3.5% NaCl solution at E = E, but only at N = 1400 cycles in air. The effect of the corrosive solution at E = E, is mainly associated with an acceleration of the micropropagation in surface of the type I cracks which are more numerous than in air at low values of N. Finally, the analysis of the replicas at low strain amplitudes reveals that: (i) in air, the formation of type I cracks is very delayed in comparison to the behaviour at high strain amplitude. Type I cracks begin to form at about 10% of Ni at A,t,/2 = 4.10Y3 but only at 60% of N, at AcP/2 = 10m4.Moreover the total number of cracks decreases by a factor 10 when A&,/2 decreases from 4-10-3 to 10m4 [17]. (ii) However, at E = 4, the reduction of the fatigue life in air is similar at high and low strain amplitudes (Fig. 6). This observation emphasizes: (1) the effect of time in the corrosive solution, which is much longer at low than at high strain amplitude (2) the localization of the electrochemical reaction which is more pronounced at low than at high strain amplitude because of the reduction of the number of cracks at low strain amplitude. The same kind of tests was performed in the cathodic region (for which the hydrogen reduction is pronounced). The results obtained at E = -5OOmV/SCE are described on Fig. 9: they correspond to a very marked reduction of the fatigue life. At N = 100 cycles, the number of type I cracks is identical in air and at E = - 500 mV/SCE: there is no influence in this case of the corrosive solution. But the transition from type I to type II and III cracks is considerably accelerated at E = - 500 mV/SCE. At N = N,, only one main crack is observed and the total number of cracks is very reduced in comparison to the observations in air and at E = E,,. The fatigue damage at E = -500 mV is very localized, even at high Ac,/2. The effect of the cathodic potential is mainly related to the acceleration of the micropropagation of the first type I cracks.

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and COUDREUSE:

CORROSION

Imm

FATIGUE

IN B.C.C. STAINLESS STEELS

2

15 ,^

I

JT/A

E 0

b

10

= a

5

a

‘;

1

N/Nit%)

Fig. 10. Simultaneous evolution of the microcrack length and the parameters JT and Jb during cycling at A5,/2 = 4.10-‘, t = IO-‘s-l, E = E,, (aerated solution).

These observations can be correlated to the results obtained at low plastic strain amplitudes for which the reduction of the fatigue life at E = - 500 mV is not so noticeable (Fig. 6). In fact, the analysis of the replicas at low strain amplitudes in air reveals that the formation of type I cracks in air is delayed in this case as described above. Moreover, the number of cracks is very reduced at low A.5,/2 in comparison to the behaviour at A$/2 = 4*10m3(reduction by a factor 10 at AE,/~ = 10m4). So, the acceleration of type I crack micropropagation by the hydrogen effect is delayed at low AeJ2, which can be correlated to the decrease at E = - 500 mV of the influence of the corrosive solution when the strain amplitude is reduced. Such results are quite original and interesting conclusions can be drawn concerning the effect of anodic dissolution and hydrogen reaction on the corrosion-fatigue damage mechanisms, as it will be discussed later on. 3.5. Cyclic evolution of the current transients To correlate the mechanical and the electrochemical damages, the evolution of the peak stress u, the maximum microcracks length I and the electrochemical parameters JT and Jb is simultaneously recorded on the Fig. 10 at A6,/2 =4.10-‘, i = 10-3s-1 and imposed potential E = E,. In such conditions, the cyclic strain induces dissolution resulting from a depassivation-repassivation process [3], by breakdown of the passive film at the slip bands emergence. On Fig. 10, three different stages can be observed: (i) during the first cycles which correspond to the cyclic hardening and to a multiplication of the slip lines, an increase of the current densities occurs; (ii) at the beginning of the saturation regime, a progressive decrease of JT and Jb appears, which indicates a localization of the dissolution process. When the type I cracks are formed (after 100 cycles at Ae,/2 = 4.10-‘), the decrease of JT is slowing down; (iii) when type III cracks formation is starting by micropropagation of type II cracks, Jb which reflects the overall activity of the specimen surface

N

Fig. 1I. Influence of strain rate on the evolution of Jb and AQ at A5,/2 = 4.10-‘, E = E, (aerated solution).

during strain reversals increases again (like JT) till fracture. The effect of strain rate on the corrosion fatigue life has been shown to be sensitive [13]. Figure 11 relates this phenomenon and clearly indicates the explanation of the strain rate effect. It can be seen that at E = E,, and for A~,12 = 4.10m3, the quantity of electricity generated per cycle AQ (which can be related to the mass of the dissolved metal) is larger at i = 10m3s-’ than at i = lo-* s-‘, as the reduction of the corresponding fatigue lifetimes. Finally, the influence of the applied potential E on the current transients characteristics has been analysed. When - 300 mV < E < + 500 mV/SCE, anodic peaks in tension and compression are recorded, corresponding to the breakdown of the passive layer induced by the cyclic plastic strain. In this potential range, the value of Jb and the height of the current peaks increase with the value of E, which reflects an increasing instability of the passive layer. The Fig. 12(a) and (b) described the observed peaks at N = 20 for A5,/2 = 4.10m3 and i = 10d3 s-’ at E = +500 mV [Fig. 12(a)] and E = E, [Fig. 12(b)]. When

(a)

(b)

o.ov

(cl

t

t

E=-

BCOmV

Fig. 12. Influence of the applied electrochemical potential on the current transient characteristics at A.5,/2= 4.10-) and P = IO-' s-l (aerated solution).

MAGNIN

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and COUDREUSE:

0

-500 E

CORROSION

(mV/SCE)

Fig. 13. influence of E on the fatigue lifetimes in the aerated and desaerated solutions (At,/2 = 4.10-3, c’= 10-3s-I).

E < -500 mV,SCE [Fig. 12(c)], cathodic peaks are observed essentially in tension. During cycling, the cathodic reaction is favoured and the cathodic current density JT continuously increases till the macrocrack initiation.

3.6. Behaviour in the desaerated 3.5% NaCl solution To confirm that the reduction of the fatigue life at E < - 500 mV/SCE in the aerated solution is related to an hydrogen embrittlement, tests have been performed in the desaerated solution. In such conditions, the hydrogen reduction-reaction (2H+ + 2e -- -+ H,, with H+ + em + adsorbed H + absorbed H) is favoured in regards to the oxygen reduction reaction (fOz + H,O + e- -+ 20H- ). The obtained fatigue lifetimes in the desaerated solution at different applied potential are reported on Fig. 13, for ; = IO-3 s-1. When AeJ2 =4.10-3 and E G - 500 mV/SCE, the reduction of the fatigue life is very pronounced in the desaerated solution (N, = 150 at E = -650 mV/SCE). Nevertheless, at the same applied potential, the initial current density J, (at N = 0) is quite different in the aerated and desaerated solution. Thus, tests have been performed in the two different solutions at applied potentials for which the values of J, are identical. The curves N, =f(f,,) are presented on Fig. 14. They clearly emphasize that a marked reduction of the fatigue life occurs at cathodic potential in the desaerated solution: at Jo = 2 PA/cm?, the fatigue life is reduced by a factor of 4 in the aerated solution and by a factor of 9 in the desaerated solution, Morover, it can be observed that when J0 -to (i.e. when E -+ &) the fatigue lives become similar in the two different solutions. This result confirms that the effect of hydrogen is very reduced at E = E, but particularly sensitive at E < - 500 mV/SCE.

FATIGUE

(and the dislocation behaviour) on the enhancement of the electrochemical reactions which control the reduction of the fatigue life of the b.c.c. Fe-26Cr-1 MO alloy in the NaCl30 g/l solution. This effect has been shown to be closely dependent on the imposed electrochemical potential E. The cyclic strain principally enhances the cathodic reaction at E < -500 mV/SCE and the anodic reaction at E > - 300 mV/SCE. When the cathodic reaction is favoured (E < - 500 mV/SCE) a marked reduction of the fatigue life is observed at high plastic strain amplitudes in comparison to the behaviour in air (reduction by a factor of 4-S in the aerated solution, by a factor of 9 in the desaerated solution at &,/2 = 4.10m3 and d = 10.“s-‘. The experimental data in the aerated and desaerated solutions have shown that this reduction of the fatigue life can be attributed to the enhancement of the hydrogen reduction reaction which induces the hydrogen embrittlement of the alloy. This result can be related to the well known sensitivity to hydrogen embrittlement of the ferritic stainless steels. But the effect of the cyclic strain is to favour a “local” hydrogen reduction due to the plastic deformation mode of the alloy. At high strain amplitude and for i > 10e3 so’, grain boundaries are preferential sites of the localization of the plastic strain because of the asymmetrical behaviour between tension and compression of the screw dislocations in the b.c.c. Fe-26Cr-1Mo lattice. This mechanical effect induces the intergranular type I crack initiation. The examination of the evolution of the surface damage at cathodic potentials (Fig. 9) has clearly shown that the effect of hydrogen is only sensitive when type I cracks are m~hanically formed. Then a rapid micropropagation in surface of type I cracks is observed and induces an accelerated formation of type II and type III cracks. Because of this effect the total number of cracks is much lower than in air. The corrosion fatigue damage at cathodic potential is then very localized. Such a mechanism is very relevant to explain the influence of the plastic strain amplitude on the fatigue Ni dcat.or

The experimental results presented in this paper clearly emphasize the effect of the cyclic plastic strain

on)

_-oI 100

4. DISCUSSION

2111

IN B.C.C. STAINLESS STEELS

Cothodlc

, 10

,1 ,

I 1 Jot

PA/cm’

01 0.1

Anodrc

, 1

)

Fig. 14. Influence of .TO(J, = J at N = 0) on the fatigue lifetime N, in the aerated (-) and desaerated (---) solutions (de,/2 = 4.10w3, i = IO--‘s-‘).

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MAGNIN and COUDREUSE:

CORROSION FATIGUE IN B.C.C. STAINLESS STEELS

life at cathodic potentials. Since the mechanical type I crack formation is very delayed at low strain amplitude (after 60% of Ni at Ac,/2 = 10F4), the hydrogen effect can only occur in the last part of the cycling. This explains why the reduction of the fatigue life is less noticeable at low than at high strain amplitude. Moreover, the correlation between the surface damage evolution and the current transient characteristics clearly confirm the very marked influence of microcracks on the corrosion fatigue damage by hydrogen embrittlement. When the anodic reaction is favoured {E > - 300 mV/SCE), two different cases have been encountered: (i) In the passive region (for -300 mV < E < + 500 mV/SCE), the anodic dissolution results from a depassivation-repassivation process at slip band emergence and at grain boundaries. The anodic reaction is always coupled with the cathodic reaction which can be the oxygen and/or the hydrogen reduction. But the experimental data obtained in the aerated and in the desaerated solutions indicates that the corrosion-fatigue damage in this potential range is controlled by the localized anodic dissolution. In fact (1) it has been shown that for E > - 300 mV, there is no difference in the fatigue life between the aerated and the desaerated solutions, the hydrogen reduction being favoured in the last case. (2) The previously described effect of the strain amplitude at cathodic potential is not observed. (3) The analysis of the microcracking reveals a different situation at E > -3OOmV/SCE than at lower potentials. The number of type I cracks is much higher than at cathodic potentials and their formation is enhanced by the corrosion reaction contrarily to the behaviour at cathodic potential. It can be concluded that the corrosion fatigue m~hanisms at - 300 mV < E < + 500 mV/SCE are related to a localization of the dissolution. This effect is closely dependant on the applied strain rate which governs the cyclic depassivation. The influence of strain rate on the corrosion fatigue damage of the b.c.c. stainless steels at E = E0 has been described and analyzed in detail in a previous study [13]. This influence can be summarized as follows: (I) at low strain rates (i < lW4 s-l), the dissolution characteristics are not sufficient to induce corrosion fatigue damage, fatigue lifetimes are similar in air and in the corrosive solution; (2) at high strain rates (i > 10m3s-r) the amount of dissolution is much more pronounced because of an incomplete repassivation. The localization of the dissolution is governed by the plastic deformation mechanisms due to pencil glide. The inAuence of the plastic strain amplitude is then very pronounced: corrosion-fatigue crack initiation is intergranular for A&,/2 > 10m3(because of the behaviour of screw dislocation) and transgranular for Ac;,/2 < IO-’ (for which screw dislocations are not mobile enough). Fatigue lifetimes are reduced by a factor 1.5 in the corrosive solution

at pH 6; (3) at intermediate strain rate (i N 5.10m4 to 10m3s-l), the dissolution is essentially localized at grain boundaries because of the shape change of the grain, contrarily to the behaviour at higher strain rates for which the dissolution occurs both at grain boundaries and at slip bands emergence. Thus, the effect of the corrosive solution is particularly pronounced for a critical strain rate range near 1O-3 s-l corresponding to a maximum of the localization of the dissolution. The anodic dissolution due to the depassivationrepassivation process induces a multiplication of the type I cracks at grain boundaries (Fig. 8). This is followed by a more rapid mi~opropagation of type f cracks in surface and to an accelerated formation of type II and III cracks. These effects are quite different at E > -300 mV and at cathodic potentials. Thus, the total number of cracks at N = Ni is similar in air and at E = E, but higher than at E = - 500 mV/SCE. At E = E,,, the transition from type I to type II and III cracks is emphasized by the evolution of the current densities .Jr, and JT which are quite sensitive to the modification of the kinetics of repassivation due to superficial microcracking. Finally, in the potential range -3OOmV to + 500 mV, the higher the potential, the longer the fatigue life (Fig. 4). The observation of the current transients indicates an increasing instability of the passive layer when E increases (Fig. 12). Moreover Jb increases when E increases. These observations show that the dissolution is less and less localized when the applied potential varies from E0 to + 500 mV/SCE. This explains the increase of the fatigue life when E increases. (ii) In the sensitive to pitting region at E > 500 mV/SCE, the cyclic strain enhances pitting which generally occurs only at E > 800 mV/SCE in pure corrosion. Then, the fatigue lifetimes decrease because of stress concentrations around the pits, principally at grain boundaries where the depassivation is easier and due to the shape change effect, Intergranular crack initiation is then accelerated. 5. CONCLUSIONS This study clearly shows a marked effect of the applied electrochemical potential on the corrosion fatigue damage mechanisms of b.c.c. Fe-26Cr-1Mo alloys cycled at room temperature in 3.5% NaCl solutions. Three different kinds of behaviour have been pointed out. 1. In the cathodic region for E < - 300 mV/SCE, a reduction by a factor 5-9 of the fatigue life in air is observed at high strain amplitudes. This effect is retated to hydrogen embrittlement which occurs when type I microcracks are mechanically formed. This deleterious effect is considerably reduced at low strain amplitudes (for A.+,/2 < 10e3) for which the type I crack formation is very delayed.

MAGNIN

and COUDREUSE:

CORROSION

2. In the passive region for -300 mV < E < +500 mV/SCE, the localized anodic dissolution resulting from a depassivation-repassivation process induces a reduction of the fatigue life in air. The amount of reduction is closely dependent on the imposed strain rate which controls the kinetic of depassivation. A critical strain rate range (i 2 5*10-4-10-3 s-l) is observed for which the reduction reaches a maximum (reduction by a factor 2 at any strain amplitude). The effect of the anodic dissolution is related to an increase of the number of type I microcracks, which induces an acceleration of the micropropagation and coalescence in surface of these cracks till the formation of a macrocrack. Moreover, the higher the potential, the more instable the passive layer: this leads to a less localized dissolution at high potentials and to a corresponding increase of the fatigue life.

3. In the passive region for E > 500 mV/SCE, the cyclic strain enhances the formation of pits, which induces a new decrease of the fatigue life by an effect of stress concentration around the pits. Whatever the applied potential, the experimental results presented in this work describe the corrosion fatigue-damage of the Fe-25Cr-1Mo alloy in terms of localization

of the corrosion

reactions.

This factor

is closely related to the m~crocracking process in surface and its evolution before the formation of a macrocrack which can propagate in volume. The microcracking is a key to understand the physical mechanisms by which an aqueous corrosive solution can affect the fatigue life of b.c.c. stainless steels. It can be approached

by the observation

of the evo-

FATIGUE

IN B.C.C. STAINLESS STEELS

2113

lution of the current transients during cycling. A quantitative modelization is at present the subject of a further

study.

REFERENCES H. N. Hahn and D. J. Duquette, Acta metall. 26, 279 (1978). 2. B. D. Yan, G. C. Farrindon and C. Laird. Acta metall. 33, 1533 (1985). 3. T. Magnin and L. Coudreuse, Fatigue 84, Bi~ingham (edited bv C. J. Beeversl. Vol. III. D. 1447 (19841. 4. ?. Patei,- T. Pyle and %. Rollin~,~‘Metal &i. 8, 185 (1977). 5. C. Patel, Metall. Trans. A llA, 301 (1980). 6. D. J. Duquette, Fatigue and Microstructure (edited by M. Meshii), p. 339. Am. Sot. Metals, Metals Park. Ohio, (1979): I. T. Magnin and A. Desestret, M&m. Lent. Revue M&all. 11, 641 (1983). 8. T. Pyle,‘V. RoHins and D. Howard. J. E~ectr~e~ern. sot. 122, 1445 (1975). 9. C. Pat&, Corrosion Sci. 21, 145 (1981). 10. B. D. Yan, G. C. Farrington and C. Laird, Acta metall. 33, 1593 (1985). 11. T. Magnin and J. H. Driver, Mater. Sci. Engng 39, 175 (1975). 12. T. Magnin, A. Fourdeux and J. H. Driver, f’hysica status solidi (a) 65, 301 (1981). 17 T. Magnin and L. Coudreuse, Mater. Sci. Erzgrrg72, 125

%_’(,yg>).

14. T. Magnin, J. Driver, J. Lepinoux and L. P. Kubin, Ral. P&X A&. 19,467 (1984). 15. H. Mughrabi, K. Herz and X. Stark, Inr. J. Fract. 17, 193 (1981). 16. T. Magnin, L. Coudreuse and J. M. Lardon, Scripta Metall. 19, 1487 (1985). 17. L. Coudreuse, thesis, Saint-Etienne (1986).