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Effect of thermal aging on the low cycle fatigue behavior of Z3CN20.09M cast duplex stainless steel
Materials Science & Engineering A 646 (2015) 263–271
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Effect of thermal aging on the low cycle fatigue behavior of Z3CN20.09M cast duplex stainless steel Weifeng Chen a, Fei Xue b, Yang tian a, Dunji Yu a,n, Weiwei Yu b, Xu Chen a a b
School of Chemical Engineering and Technology, Tianjin University, Tianjin 300072, China Suzhou Nuclear Power Research Institute, Suzhou 215004, China
art ic l e i nf o
a b s t r a c t
Article history: Received 8 July 2015 Received in revised form 21 August 2015 Accepted 22 August 2015 Available online 24 August 2015
Nuclear grade Z3CN20.09M cast duplex stainless steel exhibits enhanced cyclic stress response and prolonged low cycle fatigue life at room temperature after thermal aging at 400 °C for up to 6000 h. The threshold strain amplitude for the onset of secondary hardening is shifted to a lower value after thermal aging. Microstructural observations reveal that fatigue cracks tend to initiate from phase boundaries in virgin specimens, but to initiate in the ferrite phase in aged ones. Denser fatigue striations are found on the fracture surface of fatigued specimen subjected to longer thermal aging duration. These observations are explained in the context of thermal aging induced embrittlement of the ferrite phase and deformation induced martensitic phase transformation in the austenite phase. & 2015 Elsevier B.V. All rights reserved.
Keywords: Duplex stainless steel Low cycle fatigue Thermal aging
1. Introduction Owing to the high strength, good weldability, high resistance to corrosion and hot cracking, cast duplex stainless steels (CDSS), composed of a certain amount of ferrite in an austenitic matrix, are widely used for the components of the primary loop in light water reactor (LWR) nuclear power plants. The nuclear grade CDSS components including pipes, elbows, pumping and valve castings are designed for a service life of 40 years at reactor operating temperatures ranging from 280 to 320 °C. It has been recognized that CDSS can be susceptible to thermal aging-induced embrittlement during the long time service [1,2]. The influence of thermal aging embrittlement on the mechanical properties of CDSS is manifested by an increase in tensile strength [3–7], a decrease in ductility [5,7] and fracture toughness [3,4,8–10], and a raise of ductile–brittle transition temperature [4]. It has been established that the hardness of ferrite phase is significantly increased after thermal aging, whereas that of austenite phase is hardly affected [5,7,8]. These variations in the mechanical properties of CDSS after thermal aging are associated with the segregation or precipitation in the ferrite phase, which includes formation of Cr-enriched α′ phase through spinodal decomposition or through the mechanism of nucleation and growth, formation of G-phase, and precipitation of carbides and nitrides at grain boundaries [2]. In general, the embrittlement of ferrite phase by thermal aging plays the vital role in the degradation of mechanical properties of CDSS. n
Corresponding author. E-mail address: [email protected] (D. Yu).
http://dx.doi.org/10.1016/j.msea.2015.08.070 0921-5093/& 2015 Elsevier B.V. All rights reserved.
While the mechanical properties of CDSS affected by thermal aging are well documented for the monotonic loading conditions, very little data has been published concerning the influence of thermal aging on the fatigue properties of nuclear grade CDSS. In the available reports, Mager et al. [11] compared the low cycle fatigue behavior and fatigue crack growth rate of a CF8M CDSS (17.5% ferrite in volume fraction) in virgin and aged conditions, and concluded that low cycle fatigue properties and fatigue crack growth rates were not significantly modified by the thermal treatment at 400 °C for 7500 h. Kwon et al. [12] conducted low cycle fatigue tests on a CF8M CDSS (9.6% ferrite in volume fraction) aged at 430 °C for up to 3600 h, and found that the fatigue life was significantly reduced with the increase of aging time. Calonne et al. [13] investigated the fatigue crack propagation behavior of a nuclear grade CDSS containing 30% ferrite in volume fraction in virgin and aged conditions, and found a slight increase of fatigue crack growth rate after aging at 400 °C for 2400 h. Based on these limited results, no consistent conclusion can be drawn on the effect of thermal aging on the fatigue properties of nuclear grade CDSS, but it can be inferred that aging affected fatigue properties may at least depend on the ferrite content, aging temperature and time. Therefore, it is still necessary to conduct research on specified CDSSs to understand the effect of thermal aging on their fatigue properties. The objective of this study is to investigate the effect of thermal aging on the low cycle fatigue behavior of a nuclear grade Z3CN20.09M CDSS (in French Specification, similar to CF8A CDSS in ASME Specification), which is used for the primary coolant pipes in pressurized water reactor (PWR) nuclear power plants. Accelerated thermal aging experiments were conducted at 400 °C
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and low cycle fatigue tests were performed at room temperature. After the fatigue tests, scanning electron microscopy (SEM) and optical microscopy (OM) were employed in order to investigate the mechanism of the fatigue crack initiation and propagation and fracture behavior.
2. Experiments 2.1. Material The chemical composition of the Z3CN20.09M CDSS investigated in this study is listed in Table 1. Solid bars in the diameter of 18 mm were cut directly from an unused primary coolant pipe centrifugally cast by Sichuan Sanzhou SCMP Nuclear Equipment Manufacture Incorporation, then were solution treated at 1085 °C for 7.5 h, followed by water quenching. The material contains both ferrite and austenite phases in the volume fraction of 16% and 84%, respectively. Ferrite phase is island-like distributed in the austenitic matrix, as shown in Fig. 1.
2.2. Accelerated thermal aging experiments Accelerated thermal aging experiments were performed on the solid bars of Z3CN20.09M CDSS for 500, 3000 and 6000 h at 400 °C, respectively. This aging temperature is about 100 °C higher than the normal service temperature (280–320 °C) so as to accelerate the microstructural changes that may occur during service [14]. Thermal aging durations of 500, 3000 and 6000 h at 400 °C correspond to 5, 20 and 40 years at the service temperature of 320 °C, respectively, according to Arrhenius equation [15]. 2.3. Fatigue tests Virgin and aged solid bars of the steel were machined into dumb-bell shaped fatigue specimens in accordance with the specification of ASTM E606 with the 27 mm long gauge section in the diameter of 10 mm. Fully reversed strain-controlled low cycle fatigue tests were carried out on a servo-hydraulic fatigue testing machine MTS810 at room temperature in the laboratory atmosphere. The strain was measured by a knife-edge extensometer with a gauge length of 20 mm which was directly placed at equidistance from the center of the gauge section of the specimen. The strain rate of 4 10 3 s 1 was employed in all fatigue tests. Four strain amplitudes, i.e. 0.2%, 0.3%, 0.5% and 0.8%, were employed in the fatigue testing of virgin and aged specimens, respectively.
2.4. Microstructure observation Optical microscopy (OM) and scanning electron microscopy (SEM) were used to examine the surface and fracture surface of fatigued specimens so as to investigate the mechanism of deformation as well as initiation and propagation of fatigue cracks.
Fig. 1. Metallographic structure of Z3CN20.09M CDSS in virgin condition.
3. Results 3.1. Cyclic stress response Fig. 2 shows the cyclic stress response curves in tests conducted under different constant total strain amplitudes for specimens in virgin and aged conditions. A distinct and consistent feature of all curves is the “hardening plateau” during the first 100 cycles. The plateau forms as a result of the initial rapid hardening followed by gradual softening, and becomes more pronounced at larger strain amplitude. After the first 100 cycles, in both virgin and aged conditions, the material exhibits continuous and gradual softening at small strain amplitudes (Δε/2 ¼0.2% and Δε/2 ¼ 0.3%), but significant secondary hardening at large strain amplitude (Δε/ 2¼ 0.8%). At the intermediate strain amplitude of Δε/2 ¼ 0.5%, gradual softening continues after the first 100 cycles till drastic drop of loads due to macro-crack propagation in the virgin condition, whereas a short saturation is achieved before final failure in the aged conditions (500 and 3000 h); especially, in the condition aged for 6000 h, a short duration of secondary hardening is observed before drastic drop of loads. Fig. 3 shows the curves of cyclic hardening ratio with respect to fatigue life at the same strain amplitude for specimens in virgin and aged conditions, where the cyclic hardening ratio (CHR) is defined as the ratio of the stress amplitude in each cycle (Δs) to the stress amplitude in the first cycle (Δs0). At small strain amplitudes (Δε/2 ¼0.2% and Δε/2 ¼0.3%), the maximum CHR is seen hardly affected by thermal aging, whereas at the intermediate and large strain amplitudes (Δε/2 ¼0.5% and Δε/2 ¼ 0.8%), the maximum CHR is significantly enhanced in aged conditions. At all strain amplitudes, the longer thermal aging duration leads to the higher hardening ratio, although the increment is relatively small. Similarly to what is observed in Fig. 2, it is more evidently seen from Fig. 3(c) that at Δε/2 ¼0.5% after the hardening plateau, as thermal aging duration increases, the material exhibits a transition from continuous softening, to a short saturation and to secondary hardening before final failure. Collectively seen from the results given in Figs. 2 and 3, the effect of thermal aging on the cyclic stress response of Z3CN20.09M CDSS is manifested by the enhanced cyclic hardening plateau and the increased tendency to exhibit secondary hardening. The mechanisms of these phenomena are associated with the embrittlement of ferrite phase during thermal aging and the
Table 1 Chemical composition of Z3CN20.09M (wt%). Element
C
S
Si
Mn
P
Cr
Ni
Mo
Cu
Gu
N
Nb þ Ta
Ti
Fe
Content
0.035
0.008
1.13
1.10
0.023
20.14
9.09
0.083
0.072
0.054
0.026
0.01
0.03
Balance
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Fig. 2. Cyclic stress response in the fatigue tests at different strain amplitudes in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
martensitic transformation of austenite phase during fatigue cycling, which will be explained in detail in the next section.
3.2. Low cycle fatigue life Fig. 4 gives the data of fatigue life as a function of thermal aging duration at different constant strain amplitudes. It is strikingly interesting to observe that at all strain amplitudes, fatigue life increases along with the extension of thermal aging duration. This observation stands out from the results reported in published literatures. For example, Mager et al. [11] found insignificant effect of thermal aging at 400 °C for 7500 h on the low cycle fatigue life of a CF8M CDSS. In contrast, Kwon et al. [12] observed remarkable reduction of low cycle fatigue life for a CF8M CDSS thermally aged at 430 °C for up to 3600 h. In the virgin or aged condition, fatigue life is found to depend on the strain amplitude, as shown in Fig.5. According to the wellknown Coffin–Manson relation, fatigue life is controlled by the cyclic plastic strain amplitude
Δεp 2
= ε‵f (2Nf )c
where
Δεp 2
(1)
is the plastic strain amplitude, ε′f is the fatigue ductility
coefficient, and c is the fatigue ductility exponent. The Coffin– Manson plots in virgin and aged conditions are shown in Fig. 5, and the fitted values of ε′f and c are listed in Table 2.
3.3. Fatigue fracture behavior Typical fatigue fracture surface morphologies of Z3CN20.09M CDSS in virgin and aged conditions are shown in Fig. 6, where the fatigue tests were conducted at the strain amplitude of 0.5%. In all conditions, the fracture surface can be divided to three regions: (I) crack initiation region, (II) crack propagation region, and (III) final fracture region, as illustrated in Fig. 6. Fatigue cracks are found to nucleate from the surface of specimens in all cases, which is normally associated with persistent slip bands or micro defects on the specimen surface [16]. Subsequently, the cracks propagate in a radial fashion from the initiation locations to a large part of the fracture surface, forming into the region II. Finally, the specimen is fractured abruptly, resulting in the coarse fracture region as compared to the crack propagation region. On the other hand, in general, thermal aging seems to influence the area of regions II and III in this way that longer thermal duration leads to larger instable fracture region, and smaller crack propagation region correspondingly. As a matter of fact, the effect of thermal aging on the fatigue fracture behavior of Z3CN20.09M CDSS can be explicit based on the following microstructural observations. 3.3.1. Crack initiation The fractured specimens shown in Fig. 6 were cut into halves through the crack initiation region along the longitudinal direction, and the cross-sections near the fracture surface were examined by using SEM to explore the mechanism of interior crack initiation. The results are shown in Fig. 7. In the virgin condition, cracks are found along phase boundaries that are almost perpendicular to the loading direction (LD). These cracks are mainly
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Fig. 3. Cyclic hardening ratio (CHR, Δs/Δs0 ) versus fatigue life curves in virgin and aged conditions at the same strain amplitude of: (a) Δε/2¼ 0.2%, (b) Δε/2¼ 0.3%, (c) Δε/ 2¼ 0.5% and (d) Δε/2 ¼0.8%.
Fig. 4. Fatigue life as a function of thermal aging duration at different strain amplitudes.
Fig. 5. Relation between plastic strain amplitude and fatigue life in virgin and aged conditions.
induced by tensile stress. In the aged conditions, cracks are found inside the ferrite phase at angles of about 20–60o to the LD. These cracks are induced by shearing stress, and the different angles are probably due to different crystallographic orientations of ferrite grains. In brief, the above results reveal that fatigue cracks initiate from phase boundaries in a tensile fashion in the virgin condition but initiate in ferrite phase in a shearing fashion in aged conditions.
3.3.2. Crack propagation Fig. 8 shows the enlarged SEM micrographs of crack propagation regions in virgin and aged conditions shown previously in Fig. 6. After cracks initiation, the opening and closing of cracks occur alternatively during each fatigue cycle, thus leading to the cumulative propagation of fatigue cracks. This is how the fatigue striations come into being as shown in Fig. 8. Each striation corresponds to one fatigue cycle period and the width of the striation represent the crack length increment in that cycle. Thus
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Table 2 Calculated values of Coffin–Manson parameters. Aging time
c
ε′f
Virgin 500 h 3000 h 6000 h
0.584 0.610 0.606 0.487
0.319 0.736 0.804 0.275
the density of striations can reflect the fatigue crack propagation rate. It is surprising to observe that fatigue striations become denser as thermal aging duration increases. This fact indicates that thermal aging helps retard the fatigue crack propagation. Fatigue striations shown in Figs. 6 and 8 indicate the transgranular crack propagation mode in virgin and aged conditions. This is further confirmed by examining the secondary cracks on the specimen surface, as shown in Fig. 9. In all conditions, cracks propagate consecutively through austenite and ferrite phase in the direction perpendicular to the LD. 3.3.3. Final fracture Fig. 10 shows the enlarged SEM micrographs of final fracture regions in virgin and aged conditions shown previously in Fig. 6. In the virgin condition, the final fracture region is full of equiaxed dimples, as shown in Fig. 10(a), indicating the ductile fracture mode under normal stress. In Fig. 10(b), the specimen aged for 500 h exhibits fewer dimples in the final fracture region than the virgin specimen. In the conditions aged for 3000 h and 6000 h, as shown in Fig. 10(c) and (d), cleavage facets and tearing ridges are observed in the final fracture region, which are the typical features of brittle fracture mode. Therefore, the evolution of final fracture mode with the extension of thermal aging duration can be characterized as a transition from ductile fracture to brittle fracture.
267
This phenomenon is consistent with the fact observed in Fig. 6 that the final fracture region becomes larger as thermal aging duration extends. Under the ductile fracture mode, internal necking contributes to the macro deformation, promotes steady crack propagation, and inhibits instable fracture; under the brittle fracture mode, the cracks tend to propagate abruptly, resulting in the large area of instable fracture.
4. Discussion The results given in Sections 3.1 and 3.2 show a very intriguing fact that thermal aging of Z3CN20.09M CDSS does not only strengthen the material but also prolong its low cycle fatigue life. Combined with the microstructure observations in Section 3.3, possible mechanisms for this fact are discussed in this section. 4.1. Mechanisms of the enhanced cyclic hardening Before we discuss the mechanisms of the enhanced cyclic hardening induced by thermal aging, we need to first understand the consistent cyclic hardening behavior in virgin and aged conditions, i.e. the hardening plateau at all strain amplitudes and the secondary hardening at large strain amplitude. Jenčuš et al. [17] observed a similar hardening plateau in the low cycle fatigue test of a duplex stainless steel SAF 2507 (containing 66% ferrite and 34% austenite in volume fraction) at the strain amplitude of 0.8%. By using in-situ neutron diffraction, they revealed that both phases are hardened during the first few cycles, contributing to the initial hardening of bulk material, while after the initial hardening, softening occurs in austenite and ferrite is not hardened or softened any more. Although the SAF 2507 steel contains much higher content of ferrite than the Z3CN20.09M CDSS, the
Fig. 6. SEM observations of fracture surfaces of fatigued specimens in tests conducted at the strain amplitude of 0.5% in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
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Fig. 7. SEM micrographs showing crack nucleation locations in the longitudinal cross-section near the fracture part in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
Fig. 8. SEM micrographs showing fatigue striations of the specimens in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
hardening/softening mechanisms can still be used for reference, considering that the deformation mechanism could be intrinsic in austenite and ferrite. Therefore, the cyclic hardening plateau of
Z3CN20.09M CDSS can be ascribed to the initial hardening of both phases and later softening in austenite. In other words, hardening/ softening in austenite contributes to the general trend of
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Fig. 9. SEM micrographs showing fatigue crack propagation in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
hardening plateau, and hardening in ferrite raises the height of the plateau. The hardening in metallic materials has been known to result from the increased level of immobile dislocation density [18], and
the softening is associated with formation of persistent slip bands, cross-slip and other complex dislocation processes that promote the mobility of pinned dislocations or the annihilation of dislocations [19–21]. These theories can be used to explain the micro
Fig. 10. SEM micrographs showing fatigue crack propagation modes in the final fracture region in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
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Fig. 11. Metallographic structures of virgin and aged materials after fatigue testing at the strain amplitude of 0.8%: in (a) virgin condition, and (b) aged condition for 6000 h.
mechanisms of hardening plateau, while another mechanism needs to be considered to account for the secondary hardening, i.e. martensitic phase transformation in austenite, which has been well recognized to cause significant secondary hardening during cyclic deformation of austenitic stainless steels [22–26]. At a given temperature, e.g. room temperature, martensitic phase transformation is found dependent on the plastic strain amplitude for some austenitic stainless steels. For example, Baudry and Pineauan [22] found that a critical plastic strain amplitude exists for occurrence of martensitic phase transformation in an Fe18Cr6.5Ni0.19C austenitic stainless steel. Bayerlein et al. [23] claimed a threshold plastic strain amplitude of 0.3% to exceed to trigger the formation of martensite in a AISI 304L stainless steel at room temperature. In the present study, we observed distinct morphology of needle-like pileups on the surface of fatigued specimens (Δε/2 ¼0.8%) in virgin and aged conditions, as shown in Fig. 11, which confirms the formation of martensite in austenite in this steel. Moreover, based on the results shown in Figs. 2 and 3, the critical strain amplitude for the onset of martensitic phase transformation appears to be 0.5% or a higher value in the virgin condition, and shifts to a lower value between 0.3% and 0.5% in aged conditions. This change may be associated with the enhanced cyclic hardening after thermal aging that causes higher stress concentration and promotes the nucleation of martensite. It has been established that in duplex stainless steels ferrite phase can be significantly embrittled after thermal aging while austenite phase is hardly affected [2,27]. The ferrite phase is embrittled by formation of Cr-enriched α′ phase through spinodal decomposition or through the mechanism of nucleation and growth, formation of G-phase, and precipitation of carbides and nitrides at grain boundaries [2]. The embrittled ferrite exhibits enhanced strength mainly reflected in terms of increased micro hardness with the increase of aging time [5,7,8,10,28]. Specially, Wang et al. [28] reported the nano-indentation hardness of ferrite and austenite in a Z3CN20.09M CDSS as a function of thermal aging time. As aging time increases, the hardness of ferrite is significantly increased, whereas that of austenite is much less affected. They considered the Cr-enriched α′ phase precipitates in the ferrite phase through spinodal decomposition as the critical reason making the dislocation slip difficult and causing the increase of strength. Li et al. [7] observed spinodal structures and G-phase homogeneous precipitates in ferrite of a NF Z3CN20.09M CDSS after long-term thermal aging, and ascribed the increase of Vickers hardness of ferrite with aging time to the prior microstructure changes. In the present study, during cyclic loading, we can reasonably assume that aging-induced Cr-enriched α′ phase or G-phase precipitates act as effective obstacles to mobile dislocations, thus contributing to the enhanced strengthening of ferrite phase. As discussed previously, the hardening of ferrite phase helps ‘lift’ the hardening plateau. Therefore, the aging-
strengthened ferrite phase could be the main contributor to the enhanced cyclic hardening of the steel after thermal aging. At small strain amplitudes (Δε/2 ¼0.2% and Δε/2 ¼0.3%), the density of mobile dislocations is relatively low due to the small range of plastic deformation, thus the strengthening effect is not obvious. At intermediate and large strain amplitudes (Δε/2 ¼ 0.5% and Δε/ 2¼ 0.8%), a large amount of mobile dislocations can be pinned by the aging-induced precipitates, thus leading to the remarkable strengthening consequence. 4.2. Mechanisms of the prolonged fatigue life Considering the fact that Z3CN20.09M CDSS is embrittled after thermal aging, as indicated by the cleavage morphology in the final fracture region shown in Fig. 10, and confirmed by the reduced impact energy and ductile fracture to brittle fracture transition observed by Xue et al. [6], normally one would expect a reduction in low cycle fatigue life. However, the results obtained in the present study show the opposite trend. Possible mechanisms can be deduced from the microstructure observations shown in Figs. 7 and 8, as discussed below. Cracks nucleate from phase boundaries in the virgin condition, as shown in Fig. 7(a), indicating the incompatible deformation in austenite phase and ferrite phase. Baczmanski et al. [29] investigated the phase specific deformation behavior of a UR45N duplex stainless steel (50% ferrite in volume fraction) by using insitu neutron diffraction. It was found that austenite phase yields earlier than ferrite phase, resulting in the incompatible deformation at small macro strains. Under small strain cycling as in the present study, incompatible deformation induced stress concentration near phase boundaries would finally lead to the segregation of phase boundaries and further formation of micro cracks. In the aged conditions, cracks are found to initiate in ferrite phase, as shown in Fig. 7(b)–(d), which is mainly due to the embrittlement of ferrite phase by thermal aging. While strengthened by spinodal decomposition and G-phase precipitations, ferrite becomes brittle and cannot bear much plastic deformation. Thus micro cracks tend to occur in ferrite to accommodate the applied macro strain. In turn, the micro cracks in ferrite help release the stress concentration near phase boundaries. After initiation, micro cracks grow to macro cracks by coalescence, and then propagate transgranularly as shown in Fig. 9. It is very interesting to notice that compared to the segregation of phase boundaries in the virgin condition, the micro cracks in ferrite phase seem to be arrested at the phase boundaries in aged conditions. To micro cracks, phase boundary is always an effective obstacle [30]. It is possible that crystallographic orientation relations between austenite and ferrite do not fulfill Kurjumov–Sachs (K–S) relations [31], thus micro cracks cannot pass through phase
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boundaries easily [32,33]. Moreover, micro cracks induced relief of stress concentration near phase boundaries helps inhibit the crack propagation from ferrite into austenite. Furthermore, martensitic phase transformation in duplex stainless steels was found beneficial to improve the fatigue resistance by either hindering the crack propagation or reducing the cracks growth rate [34]. Collectively, cracks propagation is inhibited in aged conditions, resulting in denser fatigue striations than in the virgin condition, as shown in Fig. 8, thus leading to a prolonged fatigue life.
5. Conclusions
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
In this study, the effect of thermal aging on the low cycle fatigue behavior of nuclear grade Z3CN20.09M CDSS was investigated. Some important results and conclusions were obtained as follows: (1) In virgin and aged conditions, the secondary cyclic hardening is attributed to the martensitic phase transformation in austenite. (2) In aged conditions, the enhanced cyclic hardening response is ascribed to the strengthening of ferrite phase by spinodal decomposition and G-phase precipitations during thermal aging. (3) In aged conditions, the threshold strain amplitude for the onset of secondary hardening is shifted to a lower value, which is presumably considered as the consequence of the enhanced cyclic hardening after thermal aging that causes higher stress concentration and promotes the nucleation of martensite. (4) Low cycle fatigue life is prolonged at all strain amplitudes as thermal aging duration increases. Microstructure observations show that cracks tend to initiate from phase boundaries in virgin specimens but to initiate in ferrite in aged ones, and that the fracture surface of fatigued specimens subjected to longer thermal aging duration exhibits denser fatigue striations. (5) Based on the microstructure observations, possible mechanisms of the prolonged life are deduced as that the micro cracks initiated in ferrite due to its embrittlement by thermal aging help release incompatible deformation caused stress concentration near phase boundaries, and that the propagation of micro cracks into austenite is inhibited by phase boundaries.
Acknowledgments The authors gratefully acknowledge financial support for this work from the National Natural Science Foundation of China (No. 51435012) and Ph.D. Programs Foundation of Ministry of Education of China (No. 20130032110018). The financial support from National Science and Technology Major Project (2011ZX06004-02) is also greatly acknowledged.
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Effect of thermal aging on the low cycle fatigue behavior of Z3CN20.09M cast duplex stainless steel Weifeng Chen a, Fei Xue b, Yang tian a, Dunji Yu a,n, Weiwei Yu b, Xu Chen a a b
School of Chemical Engineering and Technology, Tianjin University, Tianjin 300072, China Suzhou Nuclear Power Research Institute, Suzhou 215004, China
art ic l e i nf o
a b s t r a c t
Article history: Received 8 July 2015 Received in revised form 21 August 2015 Accepted 22 August 2015 Available online 24 August 2015
Nuclear grade Z3CN20.09M cast duplex stainless steel exhibits enhanced cyclic stress response and prolonged low cycle fatigue life at room temperature after thermal aging at 400 °C for up to 6000 h. The threshold strain amplitude for the onset of secondary hardening is shifted to a lower value after thermal aging. Microstructural observations reveal that fatigue cracks tend to initiate from phase boundaries in virgin specimens, but to initiate in the ferrite phase in aged ones. Denser fatigue striations are found on the fracture surface of fatigued specimen subjected to longer thermal aging duration. These observations are explained in the context of thermal aging induced embrittlement of the ferrite phase and deformation induced martensitic phase transformation in the austenite phase. & 2015 Elsevier B.V. All rights reserved.
Keywords: Duplex stainless steel Low cycle fatigue Thermal aging
1. Introduction Owing to the high strength, good weldability, high resistance to corrosion and hot cracking, cast duplex stainless steels (CDSS), composed of a certain amount of ferrite in an austenitic matrix, are widely used for the components of the primary loop in light water reactor (LWR) nuclear power plants. The nuclear grade CDSS components including pipes, elbows, pumping and valve castings are designed for a service life of 40 years at reactor operating temperatures ranging from 280 to 320 °C. It has been recognized that CDSS can be susceptible to thermal aging-induced embrittlement during the long time service [1,2]. The influence of thermal aging embrittlement on the mechanical properties of CDSS is manifested by an increase in tensile strength [3–7], a decrease in ductility [5,7] and fracture toughness [3,4,8–10], and a raise of ductile–brittle transition temperature [4]. It has been established that the hardness of ferrite phase is significantly increased after thermal aging, whereas that of austenite phase is hardly affected [5,7,8]. These variations in the mechanical properties of CDSS after thermal aging are associated with the segregation or precipitation in the ferrite phase, which includes formation of Cr-enriched α′ phase through spinodal decomposition or through the mechanism of nucleation and growth, formation of G-phase, and precipitation of carbides and nitrides at grain boundaries [2]. In general, the embrittlement of ferrite phase by thermal aging plays the vital role in the degradation of mechanical properties of CDSS. n
Corresponding author. E-mail address: [email protected] (D. Yu).
http://dx.doi.org/10.1016/j.msea.2015.08.070 0921-5093/& 2015 Elsevier B.V. All rights reserved.
While the mechanical properties of CDSS affected by thermal aging are well documented for the monotonic loading conditions, very little data has been published concerning the influence of thermal aging on the fatigue properties of nuclear grade CDSS. In the available reports, Mager et al. [11] compared the low cycle fatigue behavior and fatigue crack growth rate of a CF8M CDSS (17.5% ferrite in volume fraction) in virgin and aged conditions, and concluded that low cycle fatigue properties and fatigue crack growth rates were not significantly modified by the thermal treatment at 400 °C for 7500 h. Kwon et al. [12] conducted low cycle fatigue tests on a CF8M CDSS (9.6% ferrite in volume fraction) aged at 430 °C for up to 3600 h, and found that the fatigue life was significantly reduced with the increase of aging time. Calonne et al. [13] investigated the fatigue crack propagation behavior of a nuclear grade CDSS containing 30% ferrite in volume fraction in virgin and aged conditions, and found a slight increase of fatigue crack growth rate after aging at 400 °C for 2400 h. Based on these limited results, no consistent conclusion can be drawn on the effect of thermal aging on the fatigue properties of nuclear grade CDSS, but it can be inferred that aging affected fatigue properties may at least depend on the ferrite content, aging temperature and time. Therefore, it is still necessary to conduct research on specified CDSSs to understand the effect of thermal aging on their fatigue properties. The objective of this study is to investigate the effect of thermal aging on the low cycle fatigue behavior of a nuclear grade Z3CN20.09M CDSS (in French Specification, similar to CF8A CDSS in ASME Specification), which is used for the primary coolant pipes in pressurized water reactor (PWR) nuclear power plants. Accelerated thermal aging experiments were conducted at 400 °C
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and low cycle fatigue tests were performed at room temperature. After the fatigue tests, scanning electron microscopy (SEM) and optical microscopy (OM) were employed in order to investigate the mechanism of the fatigue crack initiation and propagation and fracture behavior.
2. Experiments 2.1. Material The chemical composition of the Z3CN20.09M CDSS investigated in this study is listed in Table 1. Solid bars in the diameter of 18 mm were cut directly from an unused primary coolant pipe centrifugally cast by Sichuan Sanzhou SCMP Nuclear Equipment Manufacture Incorporation, then were solution treated at 1085 °C for 7.5 h, followed by water quenching. The material contains both ferrite and austenite phases in the volume fraction of 16% and 84%, respectively. Ferrite phase is island-like distributed in the austenitic matrix, as shown in Fig. 1.
2.2. Accelerated thermal aging experiments Accelerated thermal aging experiments were performed on the solid bars of Z3CN20.09M CDSS for 500, 3000 and 6000 h at 400 °C, respectively. This aging temperature is about 100 °C higher than the normal service temperature (280–320 °C) so as to accelerate the microstructural changes that may occur during service [14]. Thermal aging durations of 500, 3000 and 6000 h at 400 °C correspond to 5, 20 and 40 years at the service temperature of 320 °C, respectively, according to Arrhenius equation [15]. 2.3. Fatigue tests Virgin and aged solid bars of the steel were machined into dumb-bell shaped fatigue specimens in accordance with the specification of ASTM E606 with the 27 mm long gauge section in the diameter of 10 mm. Fully reversed strain-controlled low cycle fatigue tests were carried out on a servo-hydraulic fatigue testing machine MTS810 at room temperature in the laboratory atmosphere. The strain was measured by a knife-edge extensometer with a gauge length of 20 mm which was directly placed at equidistance from the center of the gauge section of the specimen. The strain rate of 4 10 3 s 1 was employed in all fatigue tests. Four strain amplitudes, i.e. 0.2%, 0.3%, 0.5% and 0.8%, were employed in the fatigue testing of virgin and aged specimens, respectively.
2.4. Microstructure observation Optical microscopy (OM) and scanning electron microscopy (SEM) were used to examine the surface and fracture surface of fatigued specimens so as to investigate the mechanism of deformation as well as initiation and propagation of fatigue cracks.
Fig. 1. Metallographic structure of Z3CN20.09M CDSS in virgin condition.
3. Results 3.1. Cyclic stress response Fig. 2 shows the cyclic stress response curves in tests conducted under different constant total strain amplitudes for specimens in virgin and aged conditions. A distinct and consistent feature of all curves is the “hardening plateau” during the first 100 cycles. The plateau forms as a result of the initial rapid hardening followed by gradual softening, and becomes more pronounced at larger strain amplitude. After the first 100 cycles, in both virgin and aged conditions, the material exhibits continuous and gradual softening at small strain amplitudes (Δε/2 ¼0.2% and Δε/2 ¼ 0.3%), but significant secondary hardening at large strain amplitude (Δε/ 2¼ 0.8%). At the intermediate strain amplitude of Δε/2 ¼ 0.5%, gradual softening continues after the first 100 cycles till drastic drop of loads due to macro-crack propagation in the virgin condition, whereas a short saturation is achieved before final failure in the aged conditions (500 and 3000 h); especially, in the condition aged for 6000 h, a short duration of secondary hardening is observed before drastic drop of loads. Fig. 3 shows the curves of cyclic hardening ratio with respect to fatigue life at the same strain amplitude for specimens in virgin and aged conditions, where the cyclic hardening ratio (CHR) is defined as the ratio of the stress amplitude in each cycle (Δs) to the stress amplitude in the first cycle (Δs0). At small strain amplitudes (Δε/2 ¼0.2% and Δε/2 ¼0.3%), the maximum CHR is seen hardly affected by thermal aging, whereas at the intermediate and large strain amplitudes (Δε/2 ¼0.5% and Δε/2 ¼ 0.8%), the maximum CHR is significantly enhanced in aged conditions. At all strain amplitudes, the longer thermal aging duration leads to the higher hardening ratio, although the increment is relatively small. Similarly to what is observed in Fig. 2, it is more evidently seen from Fig. 3(c) that at Δε/2 ¼0.5% after the hardening plateau, as thermal aging duration increases, the material exhibits a transition from continuous softening, to a short saturation and to secondary hardening before final failure. Collectively seen from the results given in Figs. 2 and 3, the effect of thermal aging on the cyclic stress response of Z3CN20.09M CDSS is manifested by the enhanced cyclic hardening plateau and the increased tendency to exhibit secondary hardening. The mechanisms of these phenomena are associated with the embrittlement of ferrite phase during thermal aging and the
Table 1 Chemical composition of Z3CN20.09M (wt%). Element
C
S
Si
Mn
P
Cr
Ni
Mo
Cu
Gu
N
Nb þ Ta
Ti
Fe
Content
0.035
0.008
1.13
1.10
0.023
20.14
9.09
0.083
0.072
0.054
0.026
0.01
0.03
Balance
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Fig. 2. Cyclic stress response in the fatigue tests at different strain amplitudes in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
martensitic transformation of austenite phase during fatigue cycling, which will be explained in detail in the next section.
3.2. Low cycle fatigue life Fig. 4 gives the data of fatigue life as a function of thermal aging duration at different constant strain amplitudes. It is strikingly interesting to observe that at all strain amplitudes, fatigue life increases along with the extension of thermal aging duration. This observation stands out from the results reported in published literatures. For example, Mager et al. [11] found insignificant effect of thermal aging at 400 °C for 7500 h on the low cycle fatigue life of a CF8M CDSS. In contrast, Kwon et al. [12] observed remarkable reduction of low cycle fatigue life for a CF8M CDSS thermally aged at 430 °C for up to 3600 h. In the virgin or aged condition, fatigue life is found to depend on the strain amplitude, as shown in Fig.5. According to the wellknown Coffin–Manson relation, fatigue life is controlled by the cyclic plastic strain amplitude
Δεp 2
= ε‵f (2Nf )c
where
Δεp 2
(1)
is the plastic strain amplitude, ε′f is the fatigue ductility
coefficient, and c is the fatigue ductility exponent. The Coffin– Manson plots in virgin and aged conditions are shown in Fig. 5, and the fitted values of ε′f and c are listed in Table 2.
3.3. Fatigue fracture behavior Typical fatigue fracture surface morphologies of Z3CN20.09M CDSS in virgin and aged conditions are shown in Fig. 6, where the fatigue tests were conducted at the strain amplitude of 0.5%. In all conditions, the fracture surface can be divided to three regions: (I) crack initiation region, (II) crack propagation region, and (III) final fracture region, as illustrated in Fig. 6. Fatigue cracks are found to nucleate from the surface of specimens in all cases, which is normally associated with persistent slip bands or micro defects on the specimen surface [16]. Subsequently, the cracks propagate in a radial fashion from the initiation locations to a large part of the fracture surface, forming into the region II. Finally, the specimen is fractured abruptly, resulting in the coarse fracture region as compared to the crack propagation region. On the other hand, in general, thermal aging seems to influence the area of regions II and III in this way that longer thermal duration leads to larger instable fracture region, and smaller crack propagation region correspondingly. As a matter of fact, the effect of thermal aging on the fatigue fracture behavior of Z3CN20.09M CDSS can be explicit based on the following microstructural observations. 3.3.1. Crack initiation The fractured specimens shown in Fig. 6 were cut into halves through the crack initiation region along the longitudinal direction, and the cross-sections near the fracture surface were examined by using SEM to explore the mechanism of interior crack initiation. The results are shown in Fig. 7. In the virgin condition, cracks are found along phase boundaries that are almost perpendicular to the loading direction (LD). These cracks are mainly
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Fig. 3. Cyclic hardening ratio (CHR, Δs/Δs0 ) versus fatigue life curves in virgin and aged conditions at the same strain amplitude of: (a) Δε/2¼ 0.2%, (b) Δε/2¼ 0.3%, (c) Δε/ 2¼ 0.5% and (d) Δε/2 ¼0.8%.
Fig. 4. Fatigue life as a function of thermal aging duration at different strain amplitudes.
Fig. 5. Relation between plastic strain amplitude and fatigue life in virgin and aged conditions.
induced by tensile stress. In the aged conditions, cracks are found inside the ferrite phase at angles of about 20–60o to the LD. These cracks are induced by shearing stress, and the different angles are probably due to different crystallographic orientations of ferrite grains. In brief, the above results reveal that fatigue cracks initiate from phase boundaries in a tensile fashion in the virgin condition but initiate in ferrite phase in a shearing fashion in aged conditions.
3.3.2. Crack propagation Fig. 8 shows the enlarged SEM micrographs of crack propagation regions in virgin and aged conditions shown previously in Fig. 6. After cracks initiation, the opening and closing of cracks occur alternatively during each fatigue cycle, thus leading to the cumulative propagation of fatigue cracks. This is how the fatigue striations come into being as shown in Fig. 8. Each striation corresponds to one fatigue cycle period and the width of the striation represent the crack length increment in that cycle. Thus
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Table 2 Calculated values of Coffin–Manson parameters. Aging time
c
ε′f
Virgin 500 h 3000 h 6000 h
0.584 0.610 0.606 0.487
0.319 0.736 0.804 0.275
the density of striations can reflect the fatigue crack propagation rate. It is surprising to observe that fatigue striations become denser as thermal aging duration increases. This fact indicates that thermal aging helps retard the fatigue crack propagation. Fatigue striations shown in Figs. 6 and 8 indicate the transgranular crack propagation mode in virgin and aged conditions. This is further confirmed by examining the secondary cracks on the specimen surface, as shown in Fig. 9. In all conditions, cracks propagate consecutively through austenite and ferrite phase in the direction perpendicular to the LD. 3.3.3. Final fracture Fig. 10 shows the enlarged SEM micrographs of final fracture regions in virgin and aged conditions shown previously in Fig. 6. In the virgin condition, the final fracture region is full of equiaxed dimples, as shown in Fig. 10(a), indicating the ductile fracture mode under normal stress. In Fig. 10(b), the specimen aged for 500 h exhibits fewer dimples in the final fracture region than the virgin specimen. In the conditions aged for 3000 h and 6000 h, as shown in Fig. 10(c) and (d), cleavage facets and tearing ridges are observed in the final fracture region, which are the typical features of brittle fracture mode. Therefore, the evolution of final fracture mode with the extension of thermal aging duration can be characterized as a transition from ductile fracture to brittle fracture.
267
This phenomenon is consistent with the fact observed in Fig. 6 that the final fracture region becomes larger as thermal aging duration extends. Under the ductile fracture mode, internal necking contributes to the macro deformation, promotes steady crack propagation, and inhibits instable fracture; under the brittle fracture mode, the cracks tend to propagate abruptly, resulting in the large area of instable fracture.
4. Discussion The results given in Sections 3.1 and 3.2 show a very intriguing fact that thermal aging of Z3CN20.09M CDSS does not only strengthen the material but also prolong its low cycle fatigue life. Combined with the microstructure observations in Section 3.3, possible mechanisms for this fact are discussed in this section. 4.1. Mechanisms of the enhanced cyclic hardening Before we discuss the mechanisms of the enhanced cyclic hardening induced by thermal aging, we need to first understand the consistent cyclic hardening behavior in virgin and aged conditions, i.e. the hardening plateau at all strain amplitudes and the secondary hardening at large strain amplitude. Jenčuš et al. [17] observed a similar hardening plateau in the low cycle fatigue test of a duplex stainless steel SAF 2507 (containing 66% ferrite and 34% austenite in volume fraction) at the strain amplitude of 0.8%. By using in-situ neutron diffraction, they revealed that both phases are hardened during the first few cycles, contributing to the initial hardening of bulk material, while after the initial hardening, softening occurs in austenite and ferrite is not hardened or softened any more. Although the SAF 2507 steel contains much higher content of ferrite than the Z3CN20.09M CDSS, the
Fig. 6. SEM observations of fracture surfaces of fatigued specimens in tests conducted at the strain amplitude of 0.5% in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
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Fig. 7. SEM micrographs showing crack nucleation locations in the longitudinal cross-section near the fracture part in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
Fig. 8. SEM micrographs showing fatigue striations of the specimens in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
hardening/softening mechanisms can still be used for reference, considering that the deformation mechanism could be intrinsic in austenite and ferrite. Therefore, the cyclic hardening plateau of
Z3CN20.09M CDSS can be ascribed to the initial hardening of both phases and later softening in austenite. In other words, hardening/ softening in austenite contributes to the general trend of
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Fig. 9. SEM micrographs showing fatigue crack propagation in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
hardening plateau, and hardening in ferrite raises the height of the plateau. The hardening in metallic materials has been known to result from the increased level of immobile dislocation density [18], and
the softening is associated with formation of persistent slip bands, cross-slip and other complex dislocation processes that promote the mobility of pinned dislocations or the annihilation of dislocations [19–21]. These theories can be used to explain the micro
Fig. 10. SEM micrographs showing fatigue crack propagation modes in the final fracture region in (a) virgin condition, and conditions aged for (b) 500 h, (c) 3000 h and (d) 6000 h.
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Fig. 11. Metallographic structures of virgin and aged materials after fatigue testing at the strain amplitude of 0.8%: in (a) virgin condition, and (b) aged condition for 6000 h.
mechanisms of hardening plateau, while another mechanism needs to be considered to account for the secondary hardening, i.e. martensitic phase transformation in austenite, which has been well recognized to cause significant secondary hardening during cyclic deformation of austenitic stainless steels [22–26]. At a given temperature, e.g. room temperature, martensitic phase transformation is found dependent on the plastic strain amplitude for some austenitic stainless steels. For example, Baudry and Pineauan [22] found that a critical plastic strain amplitude exists for occurrence of martensitic phase transformation in an Fe18Cr6.5Ni0.19C austenitic stainless steel. Bayerlein et al. [23] claimed a threshold plastic strain amplitude of 0.3% to exceed to trigger the formation of martensite in a AISI 304L stainless steel at room temperature. In the present study, we observed distinct morphology of needle-like pileups on the surface of fatigued specimens (Δε/2 ¼0.8%) in virgin and aged conditions, as shown in Fig. 11, which confirms the formation of martensite in austenite in this steel. Moreover, based on the results shown in Figs. 2 and 3, the critical strain amplitude for the onset of martensitic phase transformation appears to be 0.5% or a higher value in the virgin condition, and shifts to a lower value between 0.3% and 0.5% in aged conditions. This change may be associated with the enhanced cyclic hardening after thermal aging that causes higher stress concentration and promotes the nucleation of martensite. It has been established that in duplex stainless steels ferrite phase can be significantly embrittled after thermal aging while austenite phase is hardly affected [2,27]. The ferrite phase is embrittled by formation of Cr-enriched α′ phase through spinodal decomposition or through the mechanism of nucleation and growth, formation of G-phase, and precipitation of carbides and nitrides at grain boundaries [2]. The embrittled ferrite exhibits enhanced strength mainly reflected in terms of increased micro hardness with the increase of aging time [5,7,8,10,28]. Specially, Wang et al. [28] reported the nano-indentation hardness of ferrite and austenite in a Z3CN20.09M CDSS as a function of thermal aging time. As aging time increases, the hardness of ferrite is significantly increased, whereas that of austenite is much less affected. They considered the Cr-enriched α′ phase precipitates in the ferrite phase through spinodal decomposition as the critical reason making the dislocation slip difficult and causing the increase of strength. Li et al. [7] observed spinodal structures and G-phase homogeneous precipitates in ferrite of a NF Z3CN20.09M CDSS after long-term thermal aging, and ascribed the increase of Vickers hardness of ferrite with aging time to the prior microstructure changes. In the present study, during cyclic loading, we can reasonably assume that aging-induced Cr-enriched α′ phase or G-phase precipitates act as effective obstacles to mobile dislocations, thus contributing to the enhanced strengthening of ferrite phase. As discussed previously, the hardening of ferrite phase helps ‘lift’ the hardening plateau. Therefore, the aging-
strengthened ferrite phase could be the main contributor to the enhanced cyclic hardening of the steel after thermal aging. At small strain amplitudes (Δε/2 ¼0.2% and Δε/2 ¼0.3%), the density of mobile dislocations is relatively low due to the small range of plastic deformation, thus the strengthening effect is not obvious. At intermediate and large strain amplitudes (Δε/2 ¼ 0.5% and Δε/ 2¼ 0.8%), a large amount of mobile dislocations can be pinned by the aging-induced precipitates, thus leading to the remarkable strengthening consequence. 4.2. Mechanisms of the prolonged fatigue life Considering the fact that Z3CN20.09M CDSS is embrittled after thermal aging, as indicated by the cleavage morphology in the final fracture region shown in Fig. 10, and confirmed by the reduced impact energy and ductile fracture to brittle fracture transition observed by Xue et al. [6], normally one would expect a reduction in low cycle fatigue life. However, the results obtained in the present study show the opposite trend. Possible mechanisms can be deduced from the microstructure observations shown in Figs. 7 and 8, as discussed below. Cracks nucleate from phase boundaries in the virgin condition, as shown in Fig. 7(a), indicating the incompatible deformation in austenite phase and ferrite phase. Baczmanski et al. [29] investigated the phase specific deformation behavior of a UR45N duplex stainless steel (50% ferrite in volume fraction) by using insitu neutron diffraction. It was found that austenite phase yields earlier than ferrite phase, resulting in the incompatible deformation at small macro strains. Under small strain cycling as in the present study, incompatible deformation induced stress concentration near phase boundaries would finally lead to the segregation of phase boundaries and further formation of micro cracks. In the aged conditions, cracks are found to initiate in ferrite phase, as shown in Fig. 7(b)–(d), which is mainly due to the embrittlement of ferrite phase by thermal aging. While strengthened by spinodal decomposition and G-phase precipitations, ferrite becomes brittle and cannot bear much plastic deformation. Thus micro cracks tend to occur in ferrite to accommodate the applied macro strain. In turn, the micro cracks in ferrite help release the stress concentration near phase boundaries. After initiation, micro cracks grow to macro cracks by coalescence, and then propagate transgranularly as shown in Fig. 9. It is very interesting to notice that compared to the segregation of phase boundaries in the virgin condition, the micro cracks in ferrite phase seem to be arrested at the phase boundaries in aged conditions. To micro cracks, phase boundary is always an effective obstacle [30]. It is possible that crystallographic orientation relations between austenite and ferrite do not fulfill Kurjumov–Sachs (K–S) relations [31], thus micro cracks cannot pass through phase
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boundaries easily [32,33]. Moreover, micro cracks induced relief of stress concentration near phase boundaries helps inhibit the crack propagation from ferrite into austenite. Furthermore, martensitic phase transformation in duplex stainless steels was found beneficial to improve the fatigue resistance by either hindering the crack propagation or reducing the cracks growth rate [34]. Collectively, cracks propagation is inhibited in aged conditions, resulting in denser fatigue striations than in the virgin condition, as shown in Fig. 8, thus leading to a prolonged fatigue life.
5. Conclusions
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In this study, the effect of thermal aging on the low cycle fatigue behavior of nuclear grade Z3CN20.09M CDSS was investigated. Some important results and conclusions were obtained as follows: (1) In virgin and aged conditions, the secondary cyclic hardening is attributed to the martensitic phase transformation in austenite. (2) In aged conditions, the enhanced cyclic hardening response is ascribed to the strengthening of ferrite phase by spinodal decomposition and G-phase precipitations during thermal aging. (3) In aged conditions, the threshold strain amplitude for the onset of secondary hardening is shifted to a lower value, which is presumably considered as the consequence of the enhanced cyclic hardening after thermal aging that causes higher stress concentration and promotes the nucleation of martensite. (4) Low cycle fatigue life is prolonged at all strain amplitudes as thermal aging duration increases. Microstructure observations show that cracks tend to initiate from phase boundaries in virgin specimens but to initiate in ferrite in aged ones, and that the fracture surface of fatigued specimens subjected to longer thermal aging duration exhibits denser fatigue striations. (5) Based on the microstructure observations, possible mechanisms of the prolonged life are deduced as that the micro cracks initiated in ferrite due to its embrittlement by thermal aging help release incompatible deformation caused stress concentration near phase boundaries, and that the propagation of micro cracks into austenite is inhibited by phase boundaries.
Acknowledgments The authors gratefully acknowledge financial support for this work from the National Natural Science Foundation of China (No. 51435012) and Ph.D. Programs Foundation of Ministry of Education of China (No. 20130032110018). The financial support from National Science and Technology Major Project (2011ZX06004-02) is also greatly acknowledged.
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