Pergamon - Science Direct

~

Acta mater. Vol. 44. No. 10. pp. 3967-3978. 1996

Pergamon PII S 1359-6454(96)00045-6

Copyright t* 1996 Acta Metallurgica Inc. Published by ElsevierScience Ltd Printed in Great Britain. All rights reserved 1359-645496 515.00 + 0.00

A G I N G EFFECTS ON THE CYCLIC D E F O R M A T I O N MECHANISMS OF A DUPLEX STAINLESS STEEL L. LLANES'i', A. M A T E O , L. I T U R G O Y E N and M. A N G L A D A Departamento de Ciencia de los Materiales e Ingenieria Metalfirgica, Universitat Politecnica de Catalunya, Barcelona 08028. Spain (Received 26 October 1995: in rerised form 22 Januarr 1996)

Abstract--Aging effects on the cyclic deformation mechanisms of an AISI-329 duplex stainless steel have been studied on the basis of the cyclic hardening-softening response, cyclic stress-strain curve and substructure evolution within the individual phases. The cyclic behavior of an unaged and two aged materials shows, in terms of plastic strain amplitude (er0, three well-defined stages. In the first regime, at low ep~,no differences are observed among the response of the three materials as a consequence of the dominance of "austenitic-like" deformation mechanisms for all the materials. In the second regime, at intermediate er, the cyclic behavior of unaged material is associated with a mixed "austenitic'ferritic-like'" character, mainly due to plastic activity of both phases. On the other hand, the cyclic response of aged material within this intermediate ~p~ range is rather correlated to "'austenitic-like'" cyclic deformation mechanisms because of the intrinsic brittleness of the ferritic matrix. A third regime, at relatively large ~p~,suggests a synergetic phenomenon of dislocation activity, deformation twinning and demodulation of spinodal microstructure in ferrite that enables this phase to sustain plastic deformation. Thus. in this £v~ interval, the observed mechanical and substructural behavior within ferrite may be considered as relatively similar to that observed in unaged material at much lower stress levels: and therefore is amenable to be associated with "ferritic-like'" cyclic deformation mechanisms. Finally, based on the results presented, the prevalence of "'austenitic-like'" or "'ferritic-like'" cyclic deformation mechanisms, for a given plastic strain range, is discussed in terms of the different role played by the ferritic matrix in each material investigated. depending upon its embrittlement degree. Copyright ,~' 1996 Acta Metallurgica bw.

1. INTRODUCTION Duplex stainless steels are being increasingly used as structural materials in the oil, chemical and power industries. This is mainly due to their relatively high strength and toughness as well as high resistance to corrosion and stress corrosion. In the last 10 years, as a consequence of the often load bearing duty of these materials, considerable interest has arisen concerning their fatigue behavior, especially from a cyclic deformation viewpoint. Magnin et al. [I, 2] studied the cyclic deformation behavior of a 50-50% ferrite(ct)-austenite(y) stainless steel and claimed that it was related to the basic fatigue mechanisms of ferrite at high plastic strain amplitudes (ep~), and to those of austenite at relatively low ep~. These results are supported, within the low cycle fatigue (LCF) regime, by recent studies on the effect of nitrogen content, environmental conditions and strain amplitude on the cyclic response, crack initiation behavior and fatigue life of other duplex stainless steels [3, 4]. Studies on the substructural changes, at particular stain amplitudes, of duplex stainless steels have also been reported [5, 6]. The dislocation features observed within ferrite and austenite are qualitatively similar to those previously i'Author to whom all correspondence should be addressed.

found in single-phase ferritic and austenitic stainless steels, respectively. Moreover, the individualphase substructural activity reported, in terms of ep~ range, seems to be in good agreement with the findings of Magnin et al. Thus, the published results allow speculation on correspondence between cyclic response, structural evolution of the dislocation arrangements within each phase and cyclic deformation mechanisms. In light of the foregoing results it is clear that there has been significant progress on the understanding of the cyclic deformation behavior of duplex stainless steels in the last years. However, in order to exploit such knowledge it is necessary to examine the changes in such behavior as a response of microstructural evolution during service. Some of the applications for which stainless steels are potentially considered imply operation at temperatures of about 300°C, and therefore these materials are susceptible, during their service life, to severe aging embrittlement. In the last 15 years, aging effects on the evolution of mechanical properties of duplex stainless steels exposed to the intermediate temperature range (250-500'~C) have been studied intensively. F r o m these investigations it is well known that thermal aging promotes an increase in hardness and conventional tensile properties, e.g. yield and ultimate tensile stress, as well as a decrease in ductility, toughness and impact

3967

3968

LLANES et al.: CYCLIC DEFORMATION MECHANISMS

properties (e.g. Refs [7-10]. On the other hand. aging effects on the fatigue properties of duplex stainless steels have been studied much less extensively, and the few available reports are mainly concentrated on the fatigue crack growth aspect [11-15]. Concerning cyclic deformation, Iturgoyen et al. [11, 14] have studied the influence of aging at intermediate temperatures on the LCF behavior of duplex stainless steels. From their results for two different materials it can be observed that aging promotes: (1) a pronounced softening response over the range of e.p~studied (10-3-10--'); (2) a decrease of fatigue life at high ep~; and (3) an increase of fatigue life at the lowest ept investigated. However, a more extensive study on aging effects upon the cyclic deformation behavior of duplex stainless steel has not been reported so far. Such knowledge seems to be very important if a complete fatigue behavior assessment of real components in service is desired. On the other side, the cited cyclic softening behavior is similar to that found in high chromium ferritic stainless steels aged within the same temperature range [16, 17], and this presumes the use of previous knowledge in model, single-phase ferritic alloys as an important tool for a better comprehension of aging effects on the mechanical response of duplex stainless steels. It is the aim of this paper to examine the influence of aging on the cyclic deformation mechanisms of an AISI-329 type duplex stainless steel. To meet this objective, the cyclic stress-strain response and substructure evolution of two sets of aged specimens have been studied and compared to those of unaged material, over a wide range of ~.p~.Special attention has been addressed to investigate the individual substructural changes and associated deformation mechanisms of the constitutive phases in each case.

2. EXPERIMENTAL DETAILS 2.1. Material, aging treatments microstructural changes

and associated

The material used in this investigation was an AISI-329 type duplex stainless steel of chemical composition (weight percent) 24.6 Cr, 5.4 Ni, 0.036C, 0.072N, 1.73 Mn. 0.34Si, 0.21 Cu and balance Fe. Material was supplied by ACENOR (Spain) in the form of hot-rolled cylindrical rods of 15 mm diameter. These rods were cut and subjected to a solution treatment at 1050C for 1 h. followed by water quenching. The resultant microstructure was a mixture of slightly oriented austenite grains (38%) in a ferrite matrix. Cylindrical threaded specimens with gage diameter of 6 mm and gage length of 25 mm were machined from the 15 mm rods. To establish advanced life-service cyclic deformation changes, i.e. at temperatures close to 300°C for several years, accelerated aging at higher temperatures is often carried out. Though it is

recognized that there are limitations on the range of validity of such accelerated aging, it is an usual technique for studying aging effects in duplex stainless steels. Hence, such an approach was followed in this investigation. Two sets of specimens were heat treated at 47YC for 25 and 200 h followed by water quenching. Aging at 4 7 5 C of duplex stainless steels induces several microstructural changes within the ferritic constituent. These changes may involve: (1) decomposition of ferrite in two phases: one iron rich (~) and another chromium rich (~'), either via nucleation and growth of ~' precipitates or via spinodal decomposition; (2) intermetallic G phase precipitation in ferrite; and (3) carbide precipitation at austenite/ferrite interfaces. In the material studied here neither grain boundary carbides nor G-phase precipitation was observed during 475°C aging for times as high as 215 h. On the other hand, spinodal decomposition of ferrite was clearly evidenced through the observation of intense mottling by transmission electron microscopy (TEM). The 0.2% offset yield strength (ay), fracture ductility (&) and ultimate tensile strength (aurs) of the studied duplex steel in unaged and aged conditions are shown in Table I. The strong increase of av and O'UTSas well as the decrease ofer for aged materials are mainly attributed to the spinodal decomposition observed in ferrite. 2.2. Cyclic response and substructural examination Fatigue testing was conducted in a servohydraulic INSTRON machine (Model 1342) in laboratory air and at room temperature. Before cycling, the specimens were first hand-polished and then electropolished. All the tests were carried out under fully reversed total strain control and constant total strain rate (6 × 10-3 s-~). Strain was measured using clip-on extensometers. Several tests at constant total strain amplitude (~,) were carried out to study the hardening-softening behavior associated with each set of samples. The cyclic stress-strain response of the materials was determined by the ascending-step method. At low and intermediate strain amplitudes, cyclic stress saturation was reached in all sets of specimens. Hence, the number of cycles/step for determining the cyclic stress-strain curve (CSSC) within such a plastic strain range was taken after considering the hardening-softening response from constant strain amplitude tests. On the other hand, if tests with a large number of cycles/step had been used in aged Table 1. Tensilepropertiesof unaged(UA) and aged, at 475~Cfor 25 h (A25)and 200 h (A200),AISI-329duplexstainlesssteel studied in this investigation[15] Material

UA A25 A200

ov (MPa)

+f (%)

ot+vs (MPa)

494 752 833

38.5 31+9 22.5

697 1001 1090

3969

LLANES et al.: CYCLIC DEFORMATION MECHANISMS specimens, as in those run in unaged samples, early fracture at large strain amplitudes would have been promoted because of the pronounced softening behavior observed. This would have the effect of limiting data acquisition at higher amplitudes in ascending-step tests. Taking this problem into consideration cycling of aged specimens at large amplitude steps was conducted up to the detection of a gradual load drop of about 10%, with respect to the maximum load measured at that amplitude. Then, the strain amplitude was increased by about 15-20% of the previous step-strain amplitude. Such a sequence was followed until fracture. Thus, for purposes of this paper, it has been chosen to equate such final step-stress amplitudes, at large strain amplitudes, with saturation stress amplitudes (a~) and use them as the basis for defining the CSSC within the large ep, range. In order to follow substructure evolution and associated cyclic deformation mechanisms, a large number of specimens was sectioned for observation and analysis by TEM. Conventional methods of TEM sample preparation were employed. Thin foils were made from bulk samples cut perpendicularly to the tensile axis using a diamond wheel. To observe the substructural changes, the foil samples were examined by TEM (JEOL 1200 EXII). 3. RESULTS AND DISCUSSION 3.1. Cyclic stress-strain response

Figure 1 shows the cyclic hardening-softening characteristics of unaged and aged specimens cycled within different strain amplitude regimes. The presented curves correspond to plots of stress amplitude (a~) as a function of number of cycles at a fixed e,. Hardening in the first few cycles was always observed. At low and intermediate ep, such hardening was followed by either slight and rapid (low ep~) or gradual (intermediate ~o,) softening. At both ept ranges

Table2. Testingconditionsand correspondingsaturationor half-life stress and strain valuesfor specimensshownin Fig. 1. UA: unaged: A25 and A200: aged at 475 C for 25 and 200 h. respectively Saturation or half-life a~ Saturationor Specimen Applied ~:, (MPa) half-life ~r~ UA-2 UA-4 A25-1 UA-7 A200-1 A25-3 A200-3

2.0 3.0 3.0 4.0 4.5 6.0 6.0

x x × x × × ×

10 ' l0 ~ 10 ~ 10-" 10 ~ 10 -~ 10 ~

336 450 464 525 665 656 790

1.l 3.0 2.8 1.1 5.0 1.3 1.2

x x x x × × ×

10 -4 10 ' 10 -4 10 10 ' l 0 -~ 10 -~

a steady state of saturation was finally attained. At large %~ unaged and aged specimens experienced a continuous softening, after hardening, till fracture. This softening behavior was much more pronounced in the aged materials than in the unaged steel. In ascending-step tests aged specimens were cycled up to E, (and e~) larger than those used for constant-~, tests. For the largest amplitudes applied the first few cycles of each step were characterized by load drops accompanied by audible noises characteristic of twinning. In general, the different particularities of each response must be related to the dominance of distinct cyclic deformation mechanisms at the different strain amplitude regimes. Hardening-softening responses similar to those found in aged materials at large ~p~ have been previously reported in spinodally decomposed high Cr ferritic stainless steels subjected to LCF [16, 17]. The e, applied as well as the saturation or half-life aa and ep, values associated with the cyclic responses of unaged and aged materials shown in Fig. 1 are summarized in Table 2. CSSCs for unaged and aged materials, defined as described in the experimental section, are shown in Fig. 2. It should be noted that results for both ascending-step tests and constant amplitude tests are in good agreement for the three studied sets. All the curves show three well-distinguished stages, characterized by different cyclic strain hardening rates. While the three CSSCs seem to run together in stage I (%, < 1 0 - 4 ) , significant differences are found among

1000'

1000

800~

$

*Y 8 0 0

600

"0

10-3) a similar cyclic hardening behavior, but at different stress levels, is observed for the three materials. Changes observed among the stress-strain response of the studied materials, in terms of cyclic strain hardening rates, should be related to similitudes or discrepancies among active cyclic deformation mechanisms for each material at a given stage. On the other hand, the variance of stress values found should be rather associated with the overall yield strength increment (mainly due to ferrite's embrittlement) observed for the aged materials, as presented in Table 1. 3.2. Substructural evolution

The studies on the individual dislocation structures within ferrite and austenite for the unaged duplex stainless steel investigated here, as a function of stress-strain response, have been reported in detail elsewhere [18]. In general, it has been found that at low ep~, i.e. below 10-4, dislocation activity is concentrated in austenite, whereas no discernible substructural changes in ferrite are observed. The dislocation structures found in the austenitic phase at this low ep~ range exhibit a well-defined primary planar glide character and a heterogeneous distribution. Within the "intermediate range", i.e. at ep~ between 10 _4 and 6 x 10 -4, dislocation structure changes are seen in both the ferritic matrix and the austenitic grains. In the former these changes are more pronounced than in the latter, especially at the largest amplitudes of the interval. Ferrite substructural evolution includes development of early clusters of primary dislocations into low density and

Fig. 3. UA material: developing wall structure within a ferritic grain at large ep~,g = [200].

interconnected loop patches. Substructural development in austenite is determined by activation of new slip systems and formation of denser and more uniformly distributed planar arrays than those found at lower e¢. At high •p,, i.e. above 10-3, the dominant substructure evolution is observed inside the ferritic grains, where strain localization is enhanced through the development of vein into wall structure, e.g. Fig. 3. Changes of the dislocation features in austenite at these large e~ are described in terms of structural changes within a still planar glide rather than real substructure evolution. As an example

Fig. 4. UA material: two sets of planar arrays of dislocations in an austenitic grain at large ~p~,g = [1TI].

LLANES et al.:

CYCLIC DEFORMATION MECHANISMS

3971

Z H

'\

3 I~,E ",~-. . ,. ."

f~

r"

I' ~" •

2

t

\

_c

. ...)i~j~

.=_ 0

,

t

)

,:.-..'5

~ #

° ~

• ;-~

.

fJ f.

l

•

6

m

< t

J\

3972

LLANES

et al.:


Fig. 4 shows the most developed dislocation arrangement found in austenite in this large epl range, i.e. planar arrays consisting of dense bands and dislocation-free regions. This dislocation structure resembles that of the "matrix-persistent LiJders band (PLB)", but at an early formation stage, previously reported for other materials of planar slip character [19, 20]. Dislocation structures and their evolution in aged materials were studied, as a function of increasing strain amplitude, by examining the substructural features of several specimens tested either at constant strain amplitude or following an ascending-step sequence. Opposite to what is observed in cyclically deformed unaged material, high levels of dislocation

pattern heterogeneity from grain to grain and within single grains are found at all the strain amplitudes studied. At low er~, i.e. below l0 -4, and as expected from the cyclic response reported in Section 3.1, the substructural scenario is very close to that described for unaged material at similar ep,: dislocation arrays and pile-ups of planar character in austenite and relatively isolated primary dislocations of predominantly screw character in ferrite. At higher ep,, in stage II, earlier secondary slip activation and structures with a higher degree of agglomeration than in the unaged material, at similar ep,, are observed within the austenitic phase. A nice example of such a matrix structure is shown in Fig. 5 for a material aged for 25 h. Crossing of

Fig. 6. A25 material: slip dominated by primary screw dislocations within a ferritic grain in the intermediate ~ regime, g = [011"].

LLANES et al.:


primary and secondary slip systems is associated with higher dislocation density and a higher density of activated parallel glide planes. On the other hand, substructural changes within the "embrittled'" ferritic matrix are not observed at all. For instance, Fig. 6 corresponds to the dislocation structure of a ferritic grain in the same specimen. The mottled appearance cited in Section 2.1, as a consequence of spinodal decomposition of ferrite because of aging at 475~C, should be noted. Here, primary dislocations arranged mainly as straight long screw dislocations, often with short non-screw segments, are the unique substructural features observed. Activation of a second slip system is seen in a region adjacent to the grain boundary, presumably because of strain compatibility effects. The absence of fatigue-related dislocation structures in the ferritic matrix of aged material suggests that, contrary to the substructural features found in unaged material, this phase does not sustain any plastic deformation within this ep~ range. The synergetic effect of increasing strain amplitude and passivity of ferrite, in terms of carrying plastic deformation, results in a progressive substructural evolution within the austenitic grains. This is mainly observed as a change of slip character, from planar to wavy, of this phase. Figure 7 shows a clear example of such a transition. In this micrograph of a material aged for 200 h and cyclically loaded up to a final ep~ of 10 -~, a more advanced matrix structure has directly evolved into uncondensed cells within an austenitic grain. Gradual development of these arrangements into a well-defined cell-structure is observed at higher ~pL, for example, Fig. 8. Here, microtwins (indicated by arrows) are also visible. Such substructure evolution is similar to that previously reported in single-phase austenitic stainless steels [21-23]. It is very important to point out that: (1) a complete evolution of dislocation structures in austenite of unaged material, as

3973

Fig. 7. A200 material: uncondensed cell structure within an austenitic grain at er~ limiting stage II from stage Ill, g = [1 li]. described above, is not observed even at ep~ as large as 5 x 10-3: and (2) strain localization-related planar structures, i.e. PLBs, are not found in cyclically deformed austenitic grains of aged materials. The former observation is associated with the high dislocation activity of ferrite in the unaged material. The latter is correlated to strain homogenization within austenite, in terms of multiple slip, driven by the fact that all the imposed macroscopic strain is being sustained almost exclusively by a 38% volume of material that corresponds to austenite. At ept above 10-3, significant but very heterogeneous changes within the substructural scenario of

Fig. 8. A25 material: dislocation cells, well-developedwall structure and microtwins within an austenitic grain at high e~),g = [200].

3974

LLANES et al.:


the ferritic phase are observed. The cited heterogeneity from region to region within the gage length as well as from grain to grain and in the interior of single grains is much more pronounced than at lower ep~. Furthermore, this seems to be closely related to either the relative plastic activity of adjacent grains, both

ferritic and austenitic, or grain boundary effects. Figures 9 and 10 point out substructural features that support the above description. Figure 9 shows the dislocation structures observed after cyclic deformation up to a final ep~ of 2 x 10 -~ of a sample aged for 200 h. The imaging conditions

Figure 9 - - c o n t i n u e d on f a c i n g page.

LLANES et al.:


3975

Fig. 9. A200 material: (a) evolving loop patch/channel structure in region close to a grain boundary in a ferritic grain, g = []1"i'];(b) individual screw dislocations in two slip systems in the interior of the same ferritic grain as in (a), g = [01I]; and (c) relatively dense arrangements of entangled dislocations and stacking fault contrast in an austenitic grain adjacent to ferritic grain shown in (al and (b), g = [ITI].

are different for Figs 9(a)-(c) to show simultaneously the structures in both the ferritic and the austenitic grains. While a developing vein structure (loop patches and channels in between) is observed in a region close to the grain boundary of a ferritic grain [Fig. 9(a)], long screw dislocations are the only features found in the interior of the same grain [Fig~ 9(b)]. Very interesting is the clear mottled image in plastically "not deformed" areas [interior of the grain and lower part of Fig. 9(a)], but not in the "deformed" region [near the grain boundary in Fig. 9(a)]. Figure 9(c) corresponds to a neighbor austenitic grain and it shows dense dislocation tangles as well as stacking fault contrast, and may be considered as another example of planar to wavy slip transition with incremental strain amplitude in this phase. A substructural evolution very close to that described for the ferritic matrix has been previously reported in a spinodally decomposed F e - 2 6 C r - l M o alloy [16] in association with observed LCF softening. Hence, the pronounced softening behavior in aged materials as well as the lessening effect observed in the cyclic strain hardening rate of the corresponding CSSCs at ep~ larger than 10 -2 should be related, at least partly, to these substructural changes, i.e. demodulation of the spinodal microstructure of the ferrite phase due to cyclic deformation. Figure 10 corresponds to the 25-h aged specimen tested under ascending-step conditions, up to a resulting ep~ as large as 4 x 10-3. It indicates the existence of twinning, within the embrittled ferritic matrix, as an alternative mechanism of accommodating plastic deformation. This conforms with the observation of load drops together with the occurrence of characteristic twinning sounds at the largest values of e, applied, as mentioned before. On the

other hand, deformation twins within ferrite were found not to be a significant substructural feature during the cyclic deformation of the unaged material [18]. Thus, it is evident that deformation twinning is promoted by aging, mainly because the latter induces prevalence of low temperature behavior at large ept in ferrite. In physically describing the observed deformation twins, two aspects can be highlighted. First, deformation twins in adjoining grains are intimately

Fig. 10. A25 material: deformation twin within the embrittled ferritic matrix at high e~, g = [110].

3976


related to each other. Deformation twins in one grain seem to act as stress raisers, once they impinge at grain boundaries, such as to promote further deformation twinning within the adjoining grain. This may be taken as further evidence of the significant and specific strain compatibility interactions between adjacent grains within the aged material. Second, there is a close correlation between slip and deformation twinning at the tips of twins terminating within the crystal. The inter-relationship between slip and twinning in b.c.c, crystals has been reported and discussed by other authors before (e.g. Refs [24, 25]). The fact that these slip-twinning interactions are intimately associated with significant plastic deformation leads to a definition of them as being co-responsible for the mechanical response of aged materials observed at high epL. 3.3. Cyclic deformation mechanisms

The hardening-softening behavior, the cyclic stress-strain response and the substructural features described for unaged and aged materials in the previous sections allow us to define the dominant cyclic deformation mechanisms associated with each material at a given ep~ regime. 3.3.1. Unaged material. The correlation between cyclic behavior and substructure evolution for the unaged material studied here has been discussed in detail by the authors in a recent publication [18]. At low epj the mechanical and TEM findings identify austenite as the phase that carries a large part of the macroscopic strain. Therefore, the response observed in this first regime has been associated with the prevalence of "austenitic-like" cyclic deformation mechanisms. In the intermediate range of ept, the observation of dislocation activity in both phases induces the classification of the measured response as characterized by rather a mixed "ferritic/austeniticlike" behavior. Finally, at high ep~ the pronounced development of dislocation structures observed in ferrite, as well as the similitudes of the found mechanical response to that of single-phase ferritic stainless steels, suggest it could be described in terms of "ferritic-like" cyclic deformation mechanisms. 3.3.2. Aged material. The cyclic stress-strain response and the substructural scenario of aged materials, at low ep~, may be characterized by: (1) relatively rapid saturation; (2) slow increase of a, with rising epL; and (3) prevalence of austenite as the plastically active phase. In general, these observations are very close to those of unaged material and, therefore, it may be stated that in the AISI-329 type stainless steel studied, there are no significant changes as a consequence of aging in the fundamental cyclic deformation mechanisms within this low ~ range. At ep~between 10 -4 and 10 -3 the CSSCs of the aged materials are distinguished by a very pronounced slope change. It is well known that a similar noticeable rise of as with increasing ep~ occurs in the

cyclic stress-strain response of single-phase austenitic stainless steels [21,22,26]. Jin et al. [22] have associated the observed changes in cyclic strain hardening rate of the AISI-310 steel studied with the activation of new slip systems. In this second regime, a comparison between the particular substructural scenario of austenite in unaged and aged materials results in the observation of a transition from planar to wavy slip character. Mateo et al. [18] have explained the early activation of ferrite in unaged material due to a rapid hardening of the austenitic phase at low and intermediate ept. In aged materials such a sharp rise of as with increasing ep~is much more pronounced because of the intrinsic brittleness of spinodally decomposed ferrite. However, in the latter material, such hardening in austenite is not able to induce plastic activation of ferrite due to the noticeable brittle character of this phase. The longer the aging time, the steeper the CSSC; and therefore the more suitable it is to describe the dominant cyclic deformation mechanisms as being of clear "austenitic-like'" character. A low cyclic hardening stage, similar to that observed in the CSSC of unaged material but at higher stress levels, is found in the CSSC of aged materials at large ep~. Mateo et al. [18] have carried out a detailed TEM study on cyclically deformed unaged material and have associated the large amplitude stage with purely "'ferritic-like" cyclic deformation mechanisms. This is in complete agreement with previous results from other authors in both duplex [1-4] and single-phase ferritic steels [27-29]. In explaining the observed mechanical and substructural response of aged materials, ideas previously proposed in LCF studies on similarly decomposed high chromium ferritic stainless steels are recalled [16, 17]. Park et al. [16] have attributed the observed LCF softening response for the Fe--26Cr-lMo alloy to demodulation of the spinodal microstructure due to irreversible dislocation motion, the same mechanism as observed in f.c.c, spinodal alloys [30-32]. Anglada et al. [l 7] reported a similarly noticeable influence of cyclic deformation on the spinodal microstructure demodulation of an F e 2 8 C r - 2 M o - 4 N i - N b alloy. In discussing their results, they considered the well-known slip asymmetry of b.c.c, alloys and the fact that such asymmetry becomes higher with increasing ep~ as important factors in accounting for large cyclic deformation effects. Similar ideas may be applied to this investigation. Thus, the b.c.c, character of the ferritic matrix in duplex stainless steels is suggested to induce an even more pronounced and localized dislocation glide irreversibility than in f.c.c, alloys, and therefore causes an enhancement of both the referred demodulation and the cyclic deformation effects. Furthermore, following the results obtained here, the significance of the b.c.c, crystal structure of ferrite must be reinforced because of the intrinsic cojoint action of slip and twinning in providing an additional

LLANES et al.:


cyclic deformation mechanism in aged materials at high ~rl. The previous postulates seem to be suitable for satisfactorily explaining the results found here and support the description of the cyclic stress-strain response of aged materials at large ept as associated with a transition towards dominance of"ferritic-like" cyclic deformation mechanisms. The larger the e¢ and the shorter the aging time, the earlier (in terms of number of cycles and ~r,) the occurrence of such transition. However, in the aged material studied, the demodulation-plastic activity phenomenon is highly heterogeneous. This must be related to either orientation variability of the material grains with respect to the tensile axis or grain boundary effects in terms of strain compatibility-induced stresses or local slip--twinning interactions. Thus, the effective capability of demodulation of the spinodal microstructure because of dislocation activity, as well as twinning deformation, in a given ferritic grain seems to be strongly dependent upon its intrinsic orientation and, even more important, on the mechanical and substructural scenario in neighboring grains. In general, specific studies on these possible relationships do not exist and further research in this respect is required in order to attain a better understanding of the influence of cyclic deformation on the stability of the spinodal microstructure of ferrite in aged duplex stainless steels. 4. CONCLUSIONS This investigation is aimed at addressing the effect of aging on the cyclic deformation mechanisms of an AISI-329 type duplex stainless steel. The following conclusions can be drawn from the results and analyses: (1) The cyclic behavior of unaged and aged materials can be described, based on detailed mechanical and substructural observations, in terms of three clearly discernible stages: a low range at ep~ below 10-4, an intermediate range at Ep~between 10-4 and about 10-3, and a high range, at e¢ above 10-3. Each regime is amenable to being characterized for each material by well-defined and distinct hardeningsoftening behavior, cyclic strain hardening rate and saturation stresses, all of them completely associated with the observed substructural features. (2) The cyclic behavior of the material studied is relatively independent of aging at ep~lower than 10-'; this independence is related to the prevalence of "austenitic-like" cyclic deformation mechanisms for unaged and aged materials within this interval. (3) Aging effects on the cyclic stress-strain response are very pronounced in the "intermediate" range of Ep~.At these ep~,in aged materials the austenitic phase is still the only phase that sustains plastic deformation, because of the intrinsic brittleness of ferrite. Hence, their response is described in terms of

3977

"'austenitic-like" cyclic deformation mechanisms. The longer the aging time, the more appropriate is such a description. On the other hand, plastic activity in both phases of unaged material is observed within this intermediate range: thus, the cyclic response of the unaged material is rather correlated to a mixed "ferritic/austenitic-like" behavior. (4) At large ep~, the hardening-softening response as well as the substructural development observed in unaged and aged materials indicate a transition towards a dominance of "ferritic-like" cyclic deformation mechanisms, in terms of both slip and twinning. Such a transition is suggested to be observed in aged materials because ferrite becomes able to carry out plastic deformation through gradual demodulation of the spinodal microstructure by dislocation glide irreversibility and deformation twinning. However, the heterogeneous character of this phenomenon suggests that it is strongly dependent upon intrinsic grain orientation and the individual cyclic behavior of neighboring grains. (5) The prevalence of austenitic-like or ferritic-like cyclic deformation mechanisms for a given epL range seems to be a direct consequence of the different plastic deformability of the ferritic matrix, depending upon its embrittlement degree.

Acknowledgements--The authors wish to express their gratitude to ACENOR (Spain) for kindly supplying the duplex stainless steel used for this investigation. They greatly appreciate the critical review of the manuscript by Professor C. Laird. They are also grateful to the ECSC (European Coal and Steel Community) and the Spanish government through the CICYT (Comision Interministerial de Ciencia y Tecnologia) for financial assistance and sponsorship of this investigation under Contract No 7210-MA/940 and Grants No MT-1497. respectively.

REFERENCES

1. T. Magnin and J. M. Lardon, Mater. Sci. Engng A104, 21 (1988). 2. T. Magnin, J. M. Lardon and L. Coudreuse, in Low Cycle Fatigue (edited by H. D. Solomon, G. R. Halford, L. R. Kaisand and B. N. Leis), ASTM STP 942, p. 812. ASTM. Philadelphia, PA (1988). 3. J.-B. Vogt, A. Messai and J. Foct, in Duplex Stainless Steels '94 (edited by T. G. Gooch), Vol. 1, Paper 33. TW1, Cambridge (1994). 4. F. Perdriset, T. Magnin, T. Cassange, P. Hoch and F. Dupoiron, in Duplex Stainless Steels "94 (edited by T. G. Gooch), Vol. 3, Paper 13. TWI, Cambridge (1994). 5. Y. B. Xia and Z. G. Wang, Mater. Sci. Engng AlSl, 29 (1992). 6. J. Polfik, T. Kruml and S. Degallaix, Scripta metall. mater. 29, 1553 (1993). 7. H. D. Solomon and T. D. Devine Jr, in Duplex Stainless Steels (edited by R. A. Lula), p. 693. ASM, St Louis, MO (1983). 8. O. K. Chopra and H. M. Chung, in Environmental Degradation of Materials in Nuclear Power SystemWater Reactors (edited by G. J. Theus and J. R. Weeks),

p. 737. TMS, Warrandale, PA (1988).

3978


9. P. H. Pumphrey and K. N. Akhurst, Mater. Sci. Teehnol. 6, 211 (1990). 10. L. Iturgoyen, A. Mateo, L. Llanes and M. Anglada, in Materials .for Advanced Power Engineering (edited by D. Coutsouradis et al.), Vol. 1, p. 505. Kluwer Academic Publishers, Dordrecht, The Netherlands (1994). 11. L. Iturgoyen and M. Anglada, in Proc. Int. Conf. on Stainless Steels, Vol. 1, p. 746. ISIJ, Tokyo, Japan (1991). 12. T. J. Marrow, Ph.D. thesis, Cambridge University (1991). 13. T. J. Marrow and J. E. King, in Duplex Stainless Steel "91 (edited byJ. Charles and S. Bernhardsson), Vol. 1, p. 169. Les t~ditions de Physique, Les Ulis, France (1991). 14, L. Iturgoyen, J. Alcalfi and M. Anglada, in Fatigue 93 (edited by J.-P. Bailon and J. I. Dickson), Vol. 2, p. 669. EMAS, Warley, U.K. (1993). 15, L. Iturgoyen, Ph.D. thesis, Universitat Polit~cnica de Catalunya, Barcelona (1994). 16. K.-H. Park, J. C. LaSalle and L. H. Schwartz, Acta metall. 33, 205 (1985). 17. M. Anglada, J. Rodriguez and A. Isalgu6, Scripta. metall. 23, 1633 (1989). 18. A. Mateo, L. Llanes, L. lturgoyen and M. Anglada, Acta mater. 44, 1143 (1996).

19. C. Laird. S. Stanzl, R. De La Veaux and L. Buchinger. Mater. Sci. Engng 80, 143 (1986). 20. L. Buchinger, A. S. Cheng, S. Stanzl and C. Laird. Mater. Sci. Engng 80, 155 (1986). 21. M. Gerland, J. Mendez, P. Violan and B. Ait Saadi, Mater. Sei. Engng All8, 83 (1989). 22. N. Y. Jin, C. Zhong and X. Chen, Acta metall, mater. 38, 2141 (1990). 23. J.-B. Vogt, J. Foct, C. Regnard, G. Robert and J. Dhers, Metall. Trans. 22A, 2385 (1991). 24. S. Mahajan, Acta metall. 23, 671 (1975). 25. T. Magnin, L. Coudreuse and A. Fourdeux, Mater. Sci. Engng 63, L5 (1984). 26. C. Zhong, N. Y. Jin, X. Zhou, E. Meng and X. Chen. Acta metall, mater. 38, 2135 (1990). 27. T. Magnin and J. H. Driver, Mater. Sci. Engng 39, 175 (1979). 28. S. E. Fielding and W. M. Stobbs, J. Microscopy 130, 279 (1983). 29. M. Anglada, M. Nasarre and J. A. Planell, in Stainless Steel '87, p. 474. The Institute of Metals, Brookfield (1987). 30. M. P. Quin and L. H. Schwartz, Mater. Sci. Engng 46, 249 (1980). 31. H.-R. Sinning, Acta metall. 30, 1019 (1982). 32. M. P. Quin and L. H. Schwartz, Mater. Sei. Engng ,~4, 121 (1982).