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The tensile and low-cycle fatigue behavior of cold worked 316L stainless steel: influence of dynamic strain aging
International Journal of Fatigue 26 (2004) 899–910 www.elsevier.com/locate/ijfatigue
The tensile and low-cycle fatigue behavior of cold worked 316L stainless steel: influence of dynamic strain aging Seong-Gu Hong, Soon-Bok Lee Department of Mechanical Engineering, Korea Advanced Institute of Science and Technology, 373-1 Guseong-dong, Yuseong-gu, 305-701 Daejeon, South Korea. Received 18 March 2003; received in revised form 8 October 2003; accepted 5 December 2003
Abstract v
Tensile and low-cycle fatigue (LCF) tests were carried out in air in a wide temperature range from 20 to 750 C and strain rates of 1 104 1 102 =s to investigate the influence of strain rate on tensile and LCF properties of cold worked 316L stainless steel, especially in the dynamic strain aging (DSA) regime. Dynamic strain aging caused a change in the tensile properties such as v strength and ductility in the temperature range from 250 to 600 C. This temperature range coincided well with the regime of negative strain rate sensitivity. Cyclic stress response at all test conditions was characterized by an initial hardening during the few cycles, followed by gradual softening until failure. Temperature and strain rate dependence on cyclic softening appears to come from a change of the cyclic plastic deformation mechanism and the DSA effect. The regimes of DSA between tensile and LCF loading conditions in terms of the negative strain rate sensitivity were consistent with each other. At the elevated temperature, a drastic reduction in fatigue resistance was observed, and this is attributed to the effects of oxidation, creep, and dynamic strain aging and interactions among these factors. The regime of DSA accelerated a reduction in fatigue resistance by enhancing crack initiation and propagation. # 2004 Elsevier Ltd. All rights reserved. Keywords: Low-cycle fatigue; 316L stainless steel; Dynamic strain aging; Cyclic softening; Cold work
1. Introduction Type 316L stainless steel, which has been favored as a structural material in several high temperature components such as the primary side of liquid metal cooled fast breeder reactor (LMFBR), stands as a prominent material for the next generation power facilities. This is it offers a good combination of high temperature tensile and creep strength, corrosion resistance, coupled with enhanced resistance to sensitization and associated intergranular cracking. In LMFBR applications, the components are exposed to severe conditions such as v high temperatures (ranging from 300 to 600 C) and undergo temperature gradient induced cyclic thermal stresses as a result of repeated start-up and shut-down. Low-cycle fatigue (LCF) represents a predominant Corresponding author: Tel.: +82-42-869-3029; fax: +82-42-8693210. E-mail address: [email protected] (S.-B. Lee).
0142-1123/$ - see front matter # 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijfatigue.2003.12.002
failure mode: specific attention on LCF is needed in the design and life assessment of such components. In austenitic stainless steels [1–5], a reduction in fatigue resistance with increasing temperature has been reported. This was ascribed to a change in the cyclic plastic deformation mechanism, creep, and oxidation effects or interactions among these factors. It has also been v noticed that in the temperature range of 300–600 C, dynamic strain aging (DSA) induces a change in material properties such as strength and ductility, and thus a detailed research on influence of DSA is required in this temperature region. Dynamic strain aging is the phenomenon of interactions between diffusing solute atoms and mobile dislocations during plastic deformation. It depends on deformation rate and temperature, which governs the velocity of mobile dislocations and diffusing solute atoms. At low temperatures and high strain rates, the solute velocity is too slow compared with the dislocation velocity to cause a strain aging. In contrast, at
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high temperatures and low strain rates, a solute atmosphere can move with the dislocations, and the effects of DSA disappear. Therefore, DSA manifests itself in the range of intermediate strain rates and temperatures [6]. Dynamic strain aging is manifested by a serrated flow in the stress–strain curve, peaks or plateau in the variation of material strength with temperature, minima in the variation of ductility with temperature, and negative strain rate sensitivity (SRS). Mannan et al. [7] investigated DSA in 316 stainless steel. At low temperatures, the diffusing velocity of the interstitial solute atoms such as carbon and nitrogen is fast enough to move and age the dislocations. However, at high temperatures, cyclic plastic deformationinduced vacancies enhance the diffusion of chromium atoms, causing the aging of dislocations by the chromium atmospheres. The most common alloys showing dynamic strain aging are carbon steels, and is well established for monotonic tensile loading. However, few systematic investigations have been reported about the influence of DSA on low-cycle fatigue behavior, and the results are contradictory [1–5,8–10]. Abdel-Raouf et al. [9] examined the elevated temperature fatigue deformation characterization of Ferrovac E iron containing 0.007% carbon and reported that DSA causes a reduction in fatigue life. In contrast, in 2 1/4Cr–1Mo steel [10], an increase in fatigue life was observed in the temperature v range of 427–600 C due to an interactive solid solution hardening resulting from cyclic stress enhanced precipitation of Mo2C. In this study, tensile and low-cycle fatigue tests were v carried out in the temperature range (20–750 C) and strain rates of 1 104 1 102 =s for 17% cold worked 316L stainless steel. The test results are analyzed and discussed from the following viewpoints: (1) based on results of the tensile test, the influence of temperature and strain rate on tensile properties and tensile fracture mechanism, focused on the DSA regime, was investigated, and the temperature region where DSA operates was determined. (2) A change in cyclic softening behavior with temperature and strain rate was analyzed in terms of cyclic plastic deformation mechanism and DSA effect. (3) The DSA regimes for the tensile and fatigue loading conditions were compared with respect to the strain rate sensitivity. (4) The reasons for a reduction in fatigue resistance at elevated temperatures are discussed by considering the influence of dynamic strain aging on crack initiation and propagation. 2. Experiment 2.1. Material and specimen The material used in this study was 17% cold worked (CW) 316L stainless steel having a chemical
composition (wt%): C 0.025, Si 0.41, Mn 1.41, P 0.025, S 0.025, Ni 10.22, Cr 16.16, Mo 2.09, and N 0.043. The as-received material was manufactured by the following v process: material was solution-treated at 1100 C for 40 min, then water-quenched, and finally a round bar having a diameter of 16 mm was made through the cold drawn process which introduced a 17% prestrain. This treatment yielded an average grain size of 44.2 lm. The as-received material was fabricated into the dog-bone type specimens having a gauge length of 36 mm and a gauge diameter of 8 mm in accordance with ASTM standard E606-92. The specimen surface was polished along the longitudinal direction, using emery paper down to #2000 (13 lm), so as to remove surface defects such as machining marks. 2.2. Test equipment A closed-loop servo-hydraulic test system with 5-ton capacity was used to accomplish the tensile and fatigue tests. A three-zone resistance type furnace which conv trols temperature within a range of 1 C at steady state was used for temperature control. A high temperature extensometer (MTS model no.: 632-13F-20, gauge length: 25 mm) was used to control strain signal and measure strain. 2.3. Test procedures Tensile tests were conducted in laboratory air as a constant cross-head speed of 0.2, 2, and 20 mm/min in v the temperature range from 20 to 750 C. For displacement control (that is, cross-head control), it was impossible to measure the exact value of strain along the gauge length of a specimen during deformation. Thus, the displacement signal was converted to a strain signal by dividing it by the effective gauge length. An extensometer which has a travel length of 10% was used to determine the effective gauge length of the specimen. Cross-head speeds of 0.2, 2, and 20 mm/min correspond to strain rates of 1 104 , 1 103 , and 1 102 =s, respectively. Low-cycle fatigue tests were carried out in laboratory air and under fully reversed total axial strain control mode using a triangular waveform at 400, 550, v 600, and 650 C. A strain amplitude of 0.5% was used for all the tests, and the strain rates were 1 104 , 1 103 , and 1 102 =s. The displacement, load and strain signals were measured for each cycle. One cycle comprises of 200 data points. Fatigue life of the specimen was defined as a 30% load drop point of a stabilized peak load at half-life. Details of the experimental procedure can be found elsewhere [11].
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3. Results and discussion 3.1. Tensile test 3.1.1. Serrated flow in stress–strain curves The segment of stress–strain curves at a cross-head speed of 2 mm/min (_e 1 103 =s) is presented in Fig. 1. As temperature increases, the induced stress for the same strain value decreases, and Type A, B, D, E serrations are observed in the stress–strain curve. According to Mannan et al. [7], Type A, B, and C serrations have been observed for 316 stainless steel. However, in the present study, Type D serration occurred at low temperatures and A, B, E serrations at high temperatures. Type A and B serrations occurred at low strain and change over to Type E serration at values of high strain. The condition of temperature and strain rate for serration is shown in Fig. 2 and is a function of temperature and strain rate. The regime of serration moved to a higher temperature with an increase in strain rate. Type D serration was first v observed at 300 C and e_ 1 104 =s, and Type A, B, v E serrations occurred at 700 C and e_ 1 102 =s. Therefore, the regime of DSA, in terms of serrated flow v in stress–strain curves, was from 300 to 700 C and depended on the strain rate. 3.1.2. Strain rate sensitivity and tensile properties The effect of temperature and strain rate on 0.2% yield strength and true ultimate tensile strength (UTS), defined as the ‘(1 þ en )(engineering UTS)’ where en represents an engineering strain at onset of necking, is shown in Fig. 3(a) and (b). The results reveal a lack of remarkable effect of strain rate on 0.2% yield strength. However, the true UTS was significantly affected by the strain rate. As shown in Fig. 3(b), the influence of
Fig. 1.
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strain rate on true UTS can be categorized into three different temperature regions: in region I, below v 250 C, there was no significant effect of strain rate. In v region II, over 250–600 C, a negative strain rate versus stress response, which means that strain hardening decreases as the strain rate increases, was observed. In v region III, above 600 C, a positive strain rate versus stress response, which means that strain hardening increases as the strain rate increases, was observed. According to van den Beukel [12], dynamic strain aging is manifested by negative SRS value, which is defined as dr=dln_e. The value of SRS, defined as Dr=Dln_e, was calculated at each temperature, and the result is presented in Fig. 3(c). The SRS value becomes negative in v the temperature range 250–600 C, and this temperature region is considered as the regime of DSA. In general, a negative SRS regime involves serration. In the present study, it was valid at low temperatures but not at high temperatures. At high temperatures, the regime of serration was higher than the negative SRS regime. When the influence of temperature on material strength (0.2% yield strength and true UTS) is examined, material strength decreases with an increase in temperature. There exists a temperature region v (250–600 C) where a reduction of strength with increasing temperature is retarded or even slightly increased (Fig. 3(a) and (b)). Such an increase of v strength from 250 to 600 C was attributed to DSA, and is quite consistent with the negative SRS regime. Dynamic strain hardening stress, defined as the difference between true UTS and 0.2% yield strength, increases in the DSA regime since DSA effect is more remarkable on UTS than on yield strength. The variation of dynamic strain hardening stress with strain rate and temperature is shown in Fig. 3(d). A plateau
Segments of the stress–strain curves from tensile tests at e_ 1 103 =s.
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Fig. 2. Regimes of occurrence of serrated flow in the stress–strain curves.
or slightly increasing regions were observed from 250 v to 600 C and correspond well with the negative SRS regime. Dynamic strain hardening stress increased with a decrease in strain rate (that is, the DSA effect increased with a decrease in strain rate), and can be explained that low strain rate reduces the velocity of mobile dislocations. Thus, slow moving dislocations can easily be aged by the solute atmosphere. The minima in the variation of ductility with temv perature was observed from 250 to 600 C and moves to a higher temperature region with an increase in strain rate (Fig. 3(e)). The temperature region of minimum ductility is consistent with the plateau in the variation of strength with temperature and the negative SRS regime. However, it shows some differences with the serration regime. The strain rate dependence on v ductility of the material was negligible below 300 C. v Above 300 C, ductility of the material increases with decreasing strain rate, and gets more accelerated as temperature increases. The DSA of 17% CW 316L stainless steel occurred v in the temperature range 250–600 C and is manifested by a plateau in the variation of strength and dynamic strain hardening stress with temperature, a minima in the variation of ductility with temperature, and a negative SRS regime. However, it shows marginal difference in the occurrence of the serration regime. 3.1.3. Tensile fracture mechanism The fracture mechanism for monotonic tensile loading conditions was studied by observing the tensile fracture surfaces in a scanning electron microscopy (SEM). Results reveal a change in tensile fracture mechanism with temperature and strain rate (Fig. 4). A classic cup and cone type fracture in the non-DSA
regime, which exhibits three zones (Fig. 4(a)): the inner flat fibrous zone where fracture initiates, an intermediate radial zone, and an outer shear-lip zone where final v fracture occurs along a direction of 45 with the loading axis. In the DSA regime, the fibrous zone size is reduced, compared with that at other temperatures (Fig. 4(b)–(d)), and it is consistent with the result of reduction in ductility in the DSA regime. In addition, the fibrous zone in the DSA regime was twisted at final fracture. 3.2. Low-cycle fatigue test 3.2.1. Cyclic stress response Cyclic softening behavior was observed through a whole life except during the initial few cycles. This is attributed to 17% tensile prestrain, which was introduced in the as-received material by cold-working. Cyclic softening has been reported in materials, which are manufactured by cold work process or have high initial dislocation density [13–16]. In general, cyclic behavior of material resulted from an interaction between the generation of dislocations during cyclic plastic deformation and the annihilation by recovery processes. Cyclic softening of the cold-worked materials which have a high initial dislocation density, occurs when the annihilation rate of dislocations is greater than their generation rate, causing a net decrease in the dislocation density, or when a rearrangement of the dislocation structure takes place, resulting in an increase in the mean free path of dislocations [13,16]. It is suggested that this mechanism is responsible for the observed softening. As shown in Fig. 5, the observed cyclic softening behavior for all test conditions is classified into three regions: in region I, cyclic hardening is observed and peak stress reaches a maximum value
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Fig. 3. The variation of tensile properties of 17% CW 316 SS with temperature and strain rate; (a) 0.2% yield strength, (b) true ultimate tensile strength, (c) strain rate sensitivity, (d) dynamic strain hardening stress, (e) elongation.
after the initial few cycles. In region II (about 90% of fatigue life), cyclic softening occurs and continues up
until macrocrack initiation point where a steep decrement in peak stress occurs. In region III, the steep
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Fig. 4.
SEM micrographs of tensile fractured specimens at e_ 1 104 =s; (a) 200 C, (b) 300 C, (c) 400 C, (d) 550 C.
reduction in peak stress occurs because the effective load-bearing area is reduced due to the propagation of macrocracks. Static recovery, which is introduced during the thermal stabilizing process of a specimen before tests [11], is thought to be the reason for the occurrence of cyclic hardening in region I. Cyclic stress response of 17% CW 316L stainless steel is a function of temperature and can be categorized into two distinct regions. A negative strain rate v stress response in 400–600 C, transition behavior at v 600 C, and positive strain rate stress response above v 600 C. Such a tendency of cyclic stress response is consistent with that in the monotonic tensile loading conditions. The values of SRS were calculated at each temperature using the stress–strain relations at half-life (Fig. 6(a)) and compared with the tensile loading conditions (Fig. 3(c)). The results reveal a negative SRS regime (DSA regime) between tensile and LCF loading conditions coincides with each other in the temperature v range from 400 to 600 C, and SRS values under the two loading conditions are similar. As shown in Fig. 5, at each temperature, the SRS is maintained at a con-
v
v
v
v
stant value from the 20% of total life (N=Nf ¼ 0:2) to the initiation of region III (N=Nf ¼ 0:9) for all test conditions. This means that a material is stabilized, when SRS is maintained at a constant value. The negative strain rate versus stress response in the DSA regime appears to result from an increase in dislocation density during deformation. An increase in the dislocation density with decreasing strain rate has been reported in the DSA regime for 316L(N) stainless steel [5]. During DSA, slow moving dislocations become easily aged by the solute atmosphere. Thus, additional dislocations are generated to maintain the imposed deformation rate. This mechanism causes an increase in the dislocation density. The matrix is hardened by an increase in dislocation density, causing an increase in flow stress needed to impose the same total strain during successive cycles. Fig. 6(b) shows a change in plastic strain amplitude Dep =2, at a total strain amplitude of 0.5%, with temperature and strain rate. Dep =2 increased in the DSA v regime 400–550 C with an increase in strain rate, and v the reverse was observed at 650 C. It is opposite to
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Fig. 5.
v
v
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v
v
Cyclic stress responses of 17% CW 316L SS at Det ¼ 0:5%; (a) 400 C, (b) 550 C, (c) 600 C, (d) 650 C.
strain rate dependence on peak stress (Fig. 6(a)) since plastic strain amplitude is calculated from the difference between total strain amplitude and elastic strain amplitude like Eq. (1): Dep Det Dee Det Dr ¼ ¼ 2E 2 2 2 2
increases with a decrease in strain rate. This is consistent with the fact that fatigue resistance reduces with decreasing strain rate. However, temperature dependence of plastic strain energy density is ambiguous and exhibits no consistent tendency.
ð1Þ
where Dr=2 and E represent stress amplitude and elasv tic modulus. A lower Dep =2 was induced at 400–550 C due to an increase in strength by DSA, and a v maximum value is achieved at 650 C. According to Morrow [17], the plastic strain energy per cycle can be regarded as a composite measure of the amount of fatigue damage per cycle since cyclic plastic strain is related to the movement of dislocation, and cyclic stress is related to the resistance to their motion. The fatigue resistance of a metal can be characterized in terms of its capacity to absorb and dissipate plastic strain energy. Plastic strain energy density per cycle (DWp) calculated at half-life is shown in Fig. 6(c). DWp increases with decreasing strain rate at each temperature; that is, the fatigue damage per cycle
3.2.2. Cyclic softening behavior As mentioned in Section 3.2.1, cyclic softening was observed at all test conditions and depended both on temperature and strain rate. The parameter ‘‘softening ratio’’ defined as Eq. (2) was introduced to quantify the amount of cyclic softening during low-cycle fatigue. Softening ratio ¼
rmax rmax jNf =2 rmax
ð2Þ
where rmax and rmax jNf =2 represent a maximum peak stress and a peak stress at half-life, respectively. The variation of softening ratio with temperature under Det ¼ 0:5% and e_ 1 103 =s is presented in Fig. 7(a). Softening ratio (that is, the amount of cyclic softening) increases with an increase in temperature v v in the range from 20 to 200 C and above 600 C.
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Fig. 6. The variation of LCF properties of 17% CW 316L SS at half-life with temperature and strain rate at Det ¼ 0:5%; (a) peak stress, (b) the ratio between Dep and Det, (c) plastic strain energy density.
v
However, it decreases from 200 to 600 C where DSA operates. The dependence of softening ratio on temperature appears to result from a change in dislocation structure, which develops during LCF deformation. It has been reported that the dislocation structure of austenitic stainless steels involving 316L(N) stainless steel changes from cell structure at temperatures below 300 v v C to a planar one between 300 and 600 C (DSA regime) and back to cell or subgrain structure beyond v 600 C [1–5]. In the DSA regime, the mechanism of interaction between mobile dislocations and the solute atmosphere restricts dislocation cross slip and favors planar slip deformation, causing an increase in flow stress needed to impose the same total strain, and thus a decrease in the amount of cyclic softening, compared with other temperatures, occurs. In the non-DSA v v regime from 20 to 200 C and above 600 C, cell or
subgrain structures develop during low-cycle fatigue deformation, and its size increases with an increase in temperature, causing an increase in the mean free path of dislocations. This results in an increase in the amount of softening with increasing temperature. Details of the effect of temperature and strain amplitude on cyclic softening behavior of 17% CW 316L stainless steel can be found elsewhere [18]. The variation of softening with strain rate in the v temperature range 400–650 C is shown in Fig. 7(b). In v the DSA regime (400–550 C), the reduced strain rate introduced a decrease in the amount of softening. This is attributed to an enhanced DSA effect at the lower strain rate. However, in the non-DSA regime above v 600 C, the opposite effect was observed since effect of DSA is excluded and thermally assisted recovery is v enhanced at the lower strain rate. At 600 C, a transition behavior was observed.
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Fig. 7.
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The variation of softening ratio with temperature and strain rate at Det ¼ 0:5%; (a) temperature dependency, (b) strain rate dependency.
3.2.3. Low-cycle fatigue life The influence of temperature and strain rate on fatigue life at Det ¼ 0:5% is shown in Fig. 8. As shown in Fig. 8(a), a drastic reduction in fatigue life with increasing temperature was observed at each strain rate. According to the previous research [18], there is a linear relation between the logarithm of fatigue life and test temperature at a strain rate of 1 103 =s like Eq. (3): dlogNf ðTÞ ¼ C dT
ð3Þ
where the slope C means a reduction rate of fatigue resistance with increasing temperature and a higher value of C implies a rapid reduction in fatigue resistance. Eq. (3) can also be applied to the results of strain rates of 1 102 and 1 104 =s. The predicted fatigue life using Eq. (3) is plotted as a solid line in Fig. 8(a). The slope (C) is a function of strain rate and has a
value of 0.0021, 0.00152, and 0.00158 at each strain rate. A steep reduction in fatigue resistance with an increase in temperature or decrease in strain rate can be thought to come from the influence of oxidation and creep, which becomes important at elevated temperatures. Driver et al. [19] have reported a reduction in the fatigue resistance of type 316L(N) stainless steel in laboratory air with an increase in temperature compared with that in a vacuum. This was attributed to oxidation enhanced Stage I transgranular crack initiation. The oxidation effect becomes profound at the low strain rate. Considering the homologous temv perature, 400 C corresponds to about 0.36Tm (Tm: melting temperature), and the creep effect becomes important as temperature increases. A reduced strain rate introduces an increase in tensile loading time per cycle, where harmful damage is accumulated, and total test time, which results in an increase in creep damage.
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Fig. 8. The reduction of fatigue resistance with temperature and strain rate at Det ¼ 0:5%; (a) temperature dependency, (b) strain rate dependency.
When the reduction of fatigue resistance observed in the present study is explained by the effects of oxidation and creep, fatigue resistance at each temperature should reduce with decreasing strain rate, and the reduction rate in fatigue resistance with decreasing strain rate should be more severe at the higher temperatures. However, as shown in Fig. 8(b), the opposite was observed; that is, the amount of reduction in fatigue resistance with reduced strain rate decreased with an increase in temperature. This means that there is another factor, which affects fatigue resistance besides oxidation and creep, and is the DSA effect. In the investigation on Alloy 800H [2], the effects of DSA on crack initiation life and evolution of the microcrack density with the number of cycles have been studied and compared with results in vacuum, to exclude the effect of environment on crack initiation. The results reveal that DSA leads to an overall decrease in crack
initiation life, an appreciably higher crack density, and reduction in fatigue life. In other research [3,4], it has been noticed that DSA causes a faster reduction in fatigue resistance by way of rapid crack propagation where transgranular fracture is dominant. The negative strain rate stress response developed in the DSA regime introduces a higher stress concentration at the crack tip during cyclic deformation as the strain rate decreases, causing an increase in crack growth rate. This leads to a reduction in crack propagation life. It is well known that crack propagation stage is dominant in LCF conditions of ductile materials, compared with crack initiation stage, and thus the reduced crack propagation life implies a reduction in fatigue life. Therefore, it can be inferred that the steep reduction in fatigue life with reduced strain rate, in the temperature range from v 400 to 550 C (DSA regime), compared with other temperatures, is attributed to dynamic strain aging, which accelerates the reduction of fatigue resistance. In v the non-DSA regime above 600 C, the DSA effect is weakened or eliminated, and thus a slower reduction in fatigue resistance with decreasing strain rate resulted. The evidence of DSA-induced fatigue resistance reduction is manifested in Fig. 8(a). When the reduction of fatigue resistance with increasing temperature is examined at each strain rate, a steep reduction in fatigue life is observed at the high strain rate of 1 102 =s, compared with other strain rates. It is contrary to expectation since oxidation and creep effects will become more significant at low strain rate and high temperature. From the tensile test results, the regime of DSA is a function of strain rate and temperature, and moves to a higher temperature region as strain rate increases. During LCF deformation, effect v of DSA is weakened or removed beyond 600 C at the low strain rates below 1 103 =s. However, at the high strain rate of 1 102 =s, the DSA effect is still valid v even at 650 C. Hence, a severe reduction in fatigue resistance with an increase in temperature at the high strain rate of 1 102 =s can be thought to result from the DSA effect. 3.2.4. Failure mechanism in low-cycle fatigue loading conditions Fracture surfaces of LCF failed specimens were examined using SEM to study the failure mechanism in LCF loading conditions. Crack initiation occurred on the external surface in the specimen and was enhanced in the DSA regime, causing multiple crack initiation sites on the border of the fracture surface (Fig. 9(a)). v At all test conditions (400–650 C and 1 104 1 102 =s), crack propagated inward in the transgranular mode with the formation of fatigue striations (Fig. 9(b), (d) and (e)). It has been reported in SA508 [3] that dynamic strain aging would increase the degree of
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Fig. 9. SEM micrographs of LCF failed specimens at Det ¼ 0:5%; (a) 550 C, e_ ¼ 1 102 =s, (b) higher magnification view of ‘‘A’’ in (a), (c) v 650 C, e_ ¼ 1 104 =s, (d) higher magnification view of ‘‘A’’ in (c), (e) higher magnification view of ‘‘B’’ in (c). v
inhomogeniety of deformation during LCF by solute locking of moving dislocations so that DSA enhanced the partitioning of cyclic strains into segregate regions. This localized deformation leads to multiple crack initiation sites as shown in Fig. 9(a). In the study on 316L(N) stainless steel [5], where a total strain amplitude of 0.6% was employed for LCF in air, crack
initiation is transgranular and occurs from slip bands connected to the surface. Crack propagation is transgranular at higher strain rates (> 3 104 =s) in the v temperature range of 550–600 C, but at low strain rates (< 3 104 =s), intergranular cracks are noticed v on the fracture surface. Often at 600 C at these strain rates, crack propagation is assisted by oxidation. As
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shown in Fig. 9(d) and (e), the oxidation enhanced v crack propagation is observed at 650 C and strain rate of 1 104 =s, but the crack propagation mode is transgranular with fatigue striations.
4. Conclusions Based on the results and analyses on tensile and lowcycle fatigue tests, the following conclusions are made. 1. An increase in material strength and decrease in ductility were induced by dynamic strain aging in v the temperature range from 250 to 600 C, and was consistent with the regime of dynamic strain aging manifested by a plateau in the variation of dynamic strain hardening stress with temperature and a negative strain rate sensitivity. However, it shows some difference with the regime of occurrence of the serrated flow in stress–strain curves. 2. In the regime of dynamic strain aging, the fibrous zone size reduced, compared with that in other temperatures, and it is consistent with ductility reduction in the DSA regime. 3. The negative SRS regimes (DSA regimes) between tensile and LCF loading conditions coincide with each other in the temperature range from 400 to v 600 C, and SRS values under the two loading conditions are similar. 4. Softening ratio (the amount of cyclic softening during LCF) decreased in the DSA regime, compared with that in other temperatures, and it was due to the hardening by DSA. 5. A steep reduction in fatigue resistance with decreasing strain rate in the DSA regime, compared with that in other temperatures, is attributed to the enhanced crack initiation and propagation due to the DSA effect. 6. Dynamic strain aging enhanced crack initiation, causing multiple crack initiation sites, and the crack propagation mode was transgranular with fatigue v striations at all test conditions (400–650 C and 1 104 1 102 =s). Acknowledgements This work has been supported by Computer Aided Reliability Evaluation (CARE) National Research Laboratory in Korea Advanced Institute of Science and Technology (KAIST).
References [1] Kanazawa K, Yamaguchi K, Nishijima S. Mapping of low cycle fatigue mechanisms at elevated temperatures for an austenitic stainless steel. ASTM STP 942, 1988. p. 519–30. [2] Bressers J. In: Marriott JB, Merz M, Nihoul J, Ward J, editors. High temperature alloys, their exploitable potential. Amsterdam: Elsevier Applied Science; 1987. p. 385–410. [3] Valsan M, Sastry DH, Bhanu Sankara Rao K, Mannan SL. Effect of strain rate on the high-temperature low-cycle fatigue properties of a nimonic PE-16 superalloy. Metall Trans 1995; 25A:159–71. [4] Srinivasan VS, Sandhya R, Bhanu Sankara Rao K, Mannan SL, Raghavan KS. Effect of temperature on the low cycle fatigue behavior of nitrogen alloyed type 316L stainless steel. Int J Fatigue 1991;13(6):471–8. [5] Srinivasan VS, Valsan M, Sandhya R, Bhanu Sankara Rao K, Mannan SL, Sastry DH. High temperature time-dependent low cycle fatigue behavior of a type 316L(N) stainless steel. Int J Fatigue 1999;21:11–21. [6] Rodriguez P. Encyclopedia of materials science and engineering, vol. 1 (5). NY: Pergamon Press; 1988. p. 504–8. [7] Mannan SL, Samuel KG, Rodriguez P. Trans Ind Inst Metals 1983;36:313. [8] Kim DJ. The effects of dynamic strain aging on the low-cycle fatigue behaviors at high temperature for the AISI 304L stainless steel. Ph.D. thesis, Department of Material Science, KAIST, 1988. [9] Abdel-Raouf H, Plumtree A, Topper TH. Effects of temperature and deformation rate on cyclic strength and fracture of low carbon steel. ASTM STP 519, 1973. p. 28–57. [10] Challenger KD, Miller AK, Brinkman CR. An explanation for the effects of hold periods on the elevated temperature fatigue behavior of 2 1/4Cr–1Mo steel. Trans ASME J Eng Mater Tech 1981;103:7–14. [11] Hong SG, Lee SB. Development of a new LCF life prediction model of 316L stainless steel at elevated temperature. Trans KSME A 2002;26(3):521–7. [12] van den Beukel A. On the mechanism of serrated yielding and dynamic strain aging. Acta Metall 1980;28:965–9. [13] Bhanu Sankara Rao K, Valsan M, Sandhya R, Mannan SL, Rodriguez P. An assessment of cold work effects on straincontrolled low cycle fatigue behavior of type 304 stainless steel. Metall Trans 1993;24A:913–24. [14] Plumbridge WJ, Dalski ME, Castle PJ. High strain fatigue of a type 316 stainless steel. Fract Eng Mater Struct 1980;3:177–88. [15] Ganesh Sundara Raman S, Padmanabhan KA. Effect of prior cold work on the room-temperature low cycle fatigue behavior of AISI 304LN stainless steel. Int J Fatigue 1996;18(2):71–9. [16] Sherman AM. Fatigue properties of high strength-low alloy steels. Metall Trans 1975;6A:1035–40. [17] Morrow J. ASTM STP 378, 1964. p. 45–84. [18] Hong SG, Samson Y, Lee SB. The effect of temperature on lowcycle fatigue behavior of prior cold worked 316L stainless steel. Int J Fatigue 2003;25(9-11):1293–300. [19] Driver JH, Gorier C, Belrami C, Vidan P, Amzallag C. Influence of temperature and environment on the fatigue mechanisms of single-crystal and polycrystal 316L. ASTM STP 942, 1988. p. 438–55.
The tensile and low-cycle fatigue behavior of cold worked 316L stainless steel: influence of dynamic strain aging Seong-Gu Hong, Soon-Bok Lee Department of Mechanical Engineering, Korea Advanced Institute of Science and Technology, 373-1 Guseong-dong, Yuseong-gu, 305-701 Daejeon, South Korea. Received 18 March 2003; received in revised form 8 October 2003; accepted 5 December 2003
Abstract v
Tensile and low-cycle fatigue (LCF) tests were carried out in air in a wide temperature range from 20 to 750 C and strain rates of 1 104 1 102 =s to investigate the influence of strain rate on tensile and LCF properties of cold worked 316L stainless steel, especially in the dynamic strain aging (DSA) regime. Dynamic strain aging caused a change in the tensile properties such as v strength and ductility in the temperature range from 250 to 600 C. This temperature range coincided well with the regime of negative strain rate sensitivity. Cyclic stress response at all test conditions was characterized by an initial hardening during the few cycles, followed by gradual softening until failure. Temperature and strain rate dependence on cyclic softening appears to come from a change of the cyclic plastic deformation mechanism and the DSA effect. The regimes of DSA between tensile and LCF loading conditions in terms of the negative strain rate sensitivity were consistent with each other. At the elevated temperature, a drastic reduction in fatigue resistance was observed, and this is attributed to the effects of oxidation, creep, and dynamic strain aging and interactions among these factors. The regime of DSA accelerated a reduction in fatigue resistance by enhancing crack initiation and propagation. # 2004 Elsevier Ltd. All rights reserved. Keywords: Low-cycle fatigue; 316L stainless steel; Dynamic strain aging; Cyclic softening; Cold work
1. Introduction Type 316L stainless steel, which has been favored as a structural material in several high temperature components such as the primary side of liquid metal cooled fast breeder reactor (LMFBR), stands as a prominent material for the next generation power facilities. This is it offers a good combination of high temperature tensile and creep strength, corrosion resistance, coupled with enhanced resistance to sensitization and associated intergranular cracking. In LMFBR applications, the components are exposed to severe conditions such as v high temperatures (ranging from 300 to 600 C) and undergo temperature gradient induced cyclic thermal stresses as a result of repeated start-up and shut-down. Low-cycle fatigue (LCF) represents a predominant Corresponding author: Tel.: +82-42-869-3029; fax: +82-42-8693210. E-mail address: [email protected] (S.-B. Lee).
0142-1123/$ - see front matter # 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijfatigue.2003.12.002
failure mode: specific attention on LCF is needed in the design and life assessment of such components. In austenitic stainless steels [1–5], a reduction in fatigue resistance with increasing temperature has been reported. This was ascribed to a change in the cyclic plastic deformation mechanism, creep, and oxidation effects or interactions among these factors. It has also been v noticed that in the temperature range of 300–600 C, dynamic strain aging (DSA) induces a change in material properties such as strength and ductility, and thus a detailed research on influence of DSA is required in this temperature region. Dynamic strain aging is the phenomenon of interactions between diffusing solute atoms and mobile dislocations during plastic deformation. It depends on deformation rate and temperature, which governs the velocity of mobile dislocations and diffusing solute atoms. At low temperatures and high strain rates, the solute velocity is too slow compared with the dislocation velocity to cause a strain aging. In contrast, at
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high temperatures and low strain rates, a solute atmosphere can move with the dislocations, and the effects of DSA disappear. Therefore, DSA manifests itself in the range of intermediate strain rates and temperatures [6]. Dynamic strain aging is manifested by a serrated flow in the stress–strain curve, peaks or plateau in the variation of material strength with temperature, minima in the variation of ductility with temperature, and negative strain rate sensitivity (SRS). Mannan et al. [7] investigated DSA in 316 stainless steel. At low temperatures, the diffusing velocity of the interstitial solute atoms such as carbon and nitrogen is fast enough to move and age the dislocations. However, at high temperatures, cyclic plastic deformationinduced vacancies enhance the diffusion of chromium atoms, causing the aging of dislocations by the chromium atmospheres. The most common alloys showing dynamic strain aging are carbon steels, and is well established for monotonic tensile loading. However, few systematic investigations have been reported about the influence of DSA on low-cycle fatigue behavior, and the results are contradictory [1–5,8–10]. Abdel-Raouf et al. [9] examined the elevated temperature fatigue deformation characterization of Ferrovac E iron containing 0.007% carbon and reported that DSA causes a reduction in fatigue life. In contrast, in 2 1/4Cr–1Mo steel [10], an increase in fatigue life was observed in the temperature v range of 427–600 C due to an interactive solid solution hardening resulting from cyclic stress enhanced precipitation of Mo2C. In this study, tensile and low-cycle fatigue tests were v carried out in the temperature range (20–750 C) and strain rates of 1 104 1 102 =s for 17% cold worked 316L stainless steel. The test results are analyzed and discussed from the following viewpoints: (1) based on results of the tensile test, the influence of temperature and strain rate on tensile properties and tensile fracture mechanism, focused on the DSA regime, was investigated, and the temperature region where DSA operates was determined. (2) A change in cyclic softening behavior with temperature and strain rate was analyzed in terms of cyclic plastic deformation mechanism and DSA effect. (3) The DSA regimes for the tensile and fatigue loading conditions were compared with respect to the strain rate sensitivity. (4) The reasons for a reduction in fatigue resistance at elevated temperatures are discussed by considering the influence of dynamic strain aging on crack initiation and propagation. 2. Experiment 2.1. Material and specimen The material used in this study was 17% cold worked (CW) 316L stainless steel having a chemical
composition (wt%): C 0.025, Si 0.41, Mn 1.41, P 0.025, S 0.025, Ni 10.22, Cr 16.16, Mo 2.09, and N 0.043. The as-received material was manufactured by the following v process: material was solution-treated at 1100 C for 40 min, then water-quenched, and finally a round bar having a diameter of 16 mm was made through the cold drawn process which introduced a 17% prestrain. This treatment yielded an average grain size of 44.2 lm. The as-received material was fabricated into the dog-bone type specimens having a gauge length of 36 mm and a gauge diameter of 8 mm in accordance with ASTM standard E606-92. The specimen surface was polished along the longitudinal direction, using emery paper down to #2000 (13 lm), so as to remove surface defects such as machining marks. 2.2. Test equipment A closed-loop servo-hydraulic test system with 5-ton capacity was used to accomplish the tensile and fatigue tests. A three-zone resistance type furnace which conv trols temperature within a range of 1 C at steady state was used for temperature control. A high temperature extensometer (MTS model no.: 632-13F-20, gauge length: 25 mm) was used to control strain signal and measure strain. 2.3. Test procedures Tensile tests were conducted in laboratory air as a constant cross-head speed of 0.2, 2, and 20 mm/min in v the temperature range from 20 to 750 C. For displacement control (that is, cross-head control), it was impossible to measure the exact value of strain along the gauge length of a specimen during deformation. Thus, the displacement signal was converted to a strain signal by dividing it by the effective gauge length. An extensometer which has a travel length of 10% was used to determine the effective gauge length of the specimen. Cross-head speeds of 0.2, 2, and 20 mm/min correspond to strain rates of 1 104 , 1 103 , and 1 102 =s, respectively. Low-cycle fatigue tests were carried out in laboratory air and under fully reversed total axial strain control mode using a triangular waveform at 400, 550, v 600, and 650 C. A strain amplitude of 0.5% was used for all the tests, and the strain rates were 1 104 , 1 103 , and 1 102 =s. The displacement, load and strain signals were measured for each cycle. One cycle comprises of 200 data points. Fatigue life of the specimen was defined as a 30% load drop point of a stabilized peak load at half-life. Details of the experimental procedure can be found elsewhere [11].
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3. Results and discussion 3.1. Tensile test 3.1.1. Serrated flow in stress–strain curves The segment of stress–strain curves at a cross-head speed of 2 mm/min (_e 1 103 =s) is presented in Fig. 1. As temperature increases, the induced stress for the same strain value decreases, and Type A, B, D, E serrations are observed in the stress–strain curve. According to Mannan et al. [7], Type A, B, and C serrations have been observed for 316 stainless steel. However, in the present study, Type D serration occurred at low temperatures and A, B, E serrations at high temperatures. Type A and B serrations occurred at low strain and change over to Type E serration at values of high strain. The condition of temperature and strain rate for serration is shown in Fig. 2 and is a function of temperature and strain rate. The regime of serration moved to a higher temperature with an increase in strain rate. Type D serration was first v observed at 300 C and e_ 1 104 =s, and Type A, B, v E serrations occurred at 700 C and e_ 1 102 =s. Therefore, the regime of DSA, in terms of serrated flow v in stress–strain curves, was from 300 to 700 C and depended on the strain rate. 3.1.2. Strain rate sensitivity and tensile properties The effect of temperature and strain rate on 0.2% yield strength and true ultimate tensile strength (UTS), defined as the ‘(1 þ en )(engineering UTS)’ where en represents an engineering strain at onset of necking, is shown in Fig. 3(a) and (b). The results reveal a lack of remarkable effect of strain rate on 0.2% yield strength. However, the true UTS was significantly affected by the strain rate. As shown in Fig. 3(b), the influence of
Fig. 1.
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strain rate on true UTS can be categorized into three different temperature regions: in region I, below v 250 C, there was no significant effect of strain rate. In v region II, over 250–600 C, a negative strain rate versus stress response, which means that strain hardening decreases as the strain rate increases, was observed. In v region III, above 600 C, a positive strain rate versus stress response, which means that strain hardening increases as the strain rate increases, was observed. According to van den Beukel [12], dynamic strain aging is manifested by negative SRS value, which is defined as dr=dln_e. The value of SRS, defined as Dr=Dln_e, was calculated at each temperature, and the result is presented in Fig. 3(c). The SRS value becomes negative in v the temperature range 250–600 C, and this temperature region is considered as the regime of DSA. In general, a negative SRS regime involves serration. In the present study, it was valid at low temperatures but not at high temperatures. At high temperatures, the regime of serration was higher than the negative SRS regime. When the influence of temperature on material strength (0.2% yield strength and true UTS) is examined, material strength decreases with an increase in temperature. There exists a temperature region v (250–600 C) where a reduction of strength with increasing temperature is retarded or even slightly increased (Fig. 3(a) and (b)). Such an increase of v strength from 250 to 600 C was attributed to DSA, and is quite consistent with the negative SRS regime. Dynamic strain hardening stress, defined as the difference between true UTS and 0.2% yield strength, increases in the DSA regime since DSA effect is more remarkable on UTS than on yield strength. The variation of dynamic strain hardening stress with strain rate and temperature is shown in Fig. 3(d). A plateau
Segments of the stress–strain curves from tensile tests at e_ 1 103 =s.
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Fig. 2. Regimes of occurrence of serrated flow in the stress–strain curves.
or slightly increasing regions were observed from 250 v to 600 C and correspond well with the negative SRS regime. Dynamic strain hardening stress increased with a decrease in strain rate (that is, the DSA effect increased with a decrease in strain rate), and can be explained that low strain rate reduces the velocity of mobile dislocations. Thus, slow moving dislocations can easily be aged by the solute atmosphere. The minima in the variation of ductility with temv perature was observed from 250 to 600 C and moves to a higher temperature region with an increase in strain rate (Fig. 3(e)). The temperature region of minimum ductility is consistent with the plateau in the variation of strength with temperature and the negative SRS regime. However, it shows some differences with the serration regime. The strain rate dependence on v ductility of the material was negligible below 300 C. v Above 300 C, ductility of the material increases with decreasing strain rate, and gets more accelerated as temperature increases. The DSA of 17% CW 316L stainless steel occurred v in the temperature range 250–600 C and is manifested by a plateau in the variation of strength and dynamic strain hardening stress with temperature, a minima in the variation of ductility with temperature, and a negative SRS regime. However, it shows marginal difference in the occurrence of the serration regime. 3.1.3. Tensile fracture mechanism The fracture mechanism for monotonic tensile loading conditions was studied by observing the tensile fracture surfaces in a scanning electron microscopy (SEM). Results reveal a change in tensile fracture mechanism with temperature and strain rate (Fig. 4). A classic cup and cone type fracture in the non-DSA
regime, which exhibits three zones (Fig. 4(a)): the inner flat fibrous zone where fracture initiates, an intermediate radial zone, and an outer shear-lip zone where final v fracture occurs along a direction of 45 with the loading axis. In the DSA regime, the fibrous zone size is reduced, compared with that at other temperatures (Fig. 4(b)–(d)), and it is consistent with the result of reduction in ductility in the DSA regime. In addition, the fibrous zone in the DSA regime was twisted at final fracture. 3.2. Low-cycle fatigue test 3.2.1. Cyclic stress response Cyclic softening behavior was observed through a whole life except during the initial few cycles. This is attributed to 17% tensile prestrain, which was introduced in the as-received material by cold-working. Cyclic softening has been reported in materials, which are manufactured by cold work process or have high initial dislocation density [13–16]. In general, cyclic behavior of material resulted from an interaction between the generation of dislocations during cyclic plastic deformation and the annihilation by recovery processes. Cyclic softening of the cold-worked materials which have a high initial dislocation density, occurs when the annihilation rate of dislocations is greater than their generation rate, causing a net decrease in the dislocation density, or when a rearrangement of the dislocation structure takes place, resulting in an increase in the mean free path of dislocations [13,16]. It is suggested that this mechanism is responsible for the observed softening. As shown in Fig. 5, the observed cyclic softening behavior for all test conditions is classified into three regions: in region I, cyclic hardening is observed and peak stress reaches a maximum value
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Fig. 3. The variation of tensile properties of 17% CW 316 SS with temperature and strain rate; (a) 0.2% yield strength, (b) true ultimate tensile strength, (c) strain rate sensitivity, (d) dynamic strain hardening stress, (e) elongation.
after the initial few cycles. In region II (about 90% of fatigue life), cyclic softening occurs and continues up
until macrocrack initiation point where a steep decrement in peak stress occurs. In region III, the steep
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Fig. 4.
SEM micrographs of tensile fractured specimens at e_ 1 104 =s; (a) 200 C, (b) 300 C, (c) 400 C, (d) 550 C.
reduction in peak stress occurs because the effective load-bearing area is reduced due to the propagation of macrocracks. Static recovery, which is introduced during the thermal stabilizing process of a specimen before tests [11], is thought to be the reason for the occurrence of cyclic hardening in region I. Cyclic stress response of 17% CW 316L stainless steel is a function of temperature and can be categorized into two distinct regions. A negative strain rate v stress response in 400–600 C, transition behavior at v 600 C, and positive strain rate stress response above v 600 C. Such a tendency of cyclic stress response is consistent with that in the monotonic tensile loading conditions. The values of SRS were calculated at each temperature using the stress–strain relations at half-life (Fig. 6(a)) and compared with the tensile loading conditions (Fig. 3(c)). The results reveal a negative SRS regime (DSA regime) between tensile and LCF loading conditions coincides with each other in the temperature v range from 400 to 600 C, and SRS values under the two loading conditions are similar. As shown in Fig. 5, at each temperature, the SRS is maintained at a con-
v
v
v
v
stant value from the 20% of total life (N=Nf ¼ 0:2) to the initiation of region III (N=Nf ¼ 0:9) for all test conditions. This means that a material is stabilized, when SRS is maintained at a constant value. The negative strain rate versus stress response in the DSA regime appears to result from an increase in dislocation density during deformation. An increase in the dislocation density with decreasing strain rate has been reported in the DSA regime for 316L(N) stainless steel [5]. During DSA, slow moving dislocations become easily aged by the solute atmosphere. Thus, additional dislocations are generated to maintain the imposed deformation rate. This mechanism causes an increase in the dislocation density. The matrix is hardened by an increase in dislocation density, causing an increase in flow stress needed to impose the same total strain during successive cycles. Fig. 6(b) shows a change in plastic strain amplitude Dep =2, at a total strain amplitude of 0.5%, with temperature and strain rate. Dep =2 increased in the DSA v regime 400–550 C with an increase in strain rate, and v the reverse was observed at 650 C. It is opposite to
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Fig. 5.
v
v
905
v
v
Cyclic stress responses of 17% CW 316L SS at Det ¼ 0:5%; (a) 400 C, (b) 550 C, (c) 600 C, (d) 650 C.
strain rate dependence on peak stress (Fig. 6(a)) since plastic strain amplitude is calculated from the difference between total strain amplitude and elastic strain amplitude like Eq. (1): Dep Det Dee Det Dr ¼ ¼ 2E 2 2 2 2
increases with a decrease in strain rate. This is consistent with the fact that fatigue resistance reduces with decreasing strain rate. However, temperature dependence of plastic strain energy density is ambiguous and exhibits no consistent tendency.
ð1Þ
where Dr=2 and E represent stress amplitude and elasv tic modulus. A lower Dep =2 was induced at 400–550 C due to an increase in strength by DSA, and a v maximum value is achieved at 650 C. According to Morrow [17], the plastic strain energy per cycle can be regarded as a composite measure of the amount of fatigue damage per cycle since cyclic plastic strain is related to the movement of dislocation, and cyclic stress is related to the resistance to their motion. The fatigue resistance of a metal can be characterized in terms of its capacity to absorb and dissipate plastic strain energy. Plastic strain energy density per cycle (DWp) calculated at half-life is shown in Fig. 6(c). DWp increases with decreasing strain rate at each temperature; that is, the fatigue damage per cycle
3.2.2. Cyclic softening behavior As mentioned in Section 3.2.1, cyclic softening was observed at all test conditions and depended both on temperature and strain rate. The parameter ‘‘softening ratio’’ defined as Eq. (2) was introduced to quantify the amount of cyclic softening during low-cycle fatigue. Softening ratio ¼
rmax rmax jNf =2 rmax
ð2Þ
where rmax and rmax jNf =2 represent a maximum peak stress and a peak stress at half-life, respectively. The variation of softening ratio with temperature under Det ¼ 0:5% and e_ 1 103 =s is presented in Fig. 7(a). Softening ratio (that is, the amount of cyclic softening) increases with an increase in temperature v v in the range from 20 to 200 C and above 600 C.
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Fig. 6. The variation of LCF properties of 17% CW 316L SS at half-life with temperature and strain rate at Det ¼ 0:5%; (a) peak stress, (b) the ratio between Dep and Det, (c) plastic strain energy density.
v
However, it decreases from 200 to 600 C where DSA operates. The dependence of softening ratio on temperature appears to result from a change in dislocation structure, which develops during LCF deformation. It has been reported that the dislocation structure of austenitic stainless steels involving 316L(N) stainless steel changes from cell structure at temperatures below 300 v v C to a planar one between 300 and 600 C (DSA regime) and back to cell or subgrain structure beyond v 600 C [1–5]. In the DSA regime, the mechanism of interaction between mobile dislocations and the solute atmosphere restricts dislocation cross slip and favors planar slip deformation, causing an increase in flow stress needed to impose the same total strain, and thus a decrease in the amount of cyclic softening, compared with other temperatures, occurs. In the non-DSA v v regime from 20 to 200 C and above 600 C, cell or
subgrain structures develop during low-cycle fatigue deformation, and its size increases with an increase in temperature, causing an increase in the mean free path of dislocations. This results in an increase in the amount of softening with increasing temperature. Details of the effect of temperature and strain amplitude on cyclic softening behavior of 17% CW 316L stainless steel can be found elsewhere [18]. The variation of softening with strain rate in the v temperature range 400–650 C is shown in Fig. 7(b). In v the DSA regime (400–550 C), the reduced strain rate introduced a decrease in the amount of softening. This is attributed to an enhanced DSA effect at the lower strain rate. However, in the non-DSA regime above v 600 C, the opposite effect was observed since effect of DSA is excluded and thermally assisted recovery is v enhanced at the lower strain rate. At 600 C, a transition behavior was observed.
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Fig. 7.
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The variation of softening ratio with temperature and strain rate at Det ¼ 0:5%; (a) temperature dependency, (b) strain rate dependency.
3.2.3. Low-cycle fatigue life The influence of temperature and strain rate on fatigue life at Det ¼ 0:5% is shown in Fig. 8. As shown in Fig. 8(a), a drastic reduction in fatigue life with increasing temperature was observed at each strain rate. According to the previous research [18], there is a linear relation between the logarithm of fatigue life and test temperature at a strain rate of 1 103 =s like Eq. (3): dlogNf ðTÞ ¼ C dT
ð3Þ
where the slope C means a reduction rate of fatigue resistance with increasing temperature and a higher value of C implies a rapid reduction in fatigue resistance. Eq. (3) can also be applied to the results of strain rates of 1 102 and 1 104 =s. The predicted fatigue life using Eq. (3) is plotted as a solid line in Fig. 8(a). The slope (C) is a function of strain rate and has a
value of 0.0021, 0.00152, and 0.00158 at each strain rate. A steep reduction in fatigue resistance with an increase in temperature or decrease in strain rate can be thought to come from the influence of oxidation and creep, which becomes important at elevated temperatures. Driver et al. [19] have reported a reduction in the fatigue resistance of type 316L(N) stainless steel in laboratory air with an increase in temperature compared with that in a vacuum. This was attributed to oxidation enhanced Stage I transgranular crack initiation. The oxidation effect becomes profound at the low strain rate. Considering the homologous temv perature, 400 C corresponds to about 0.36Tm (Tm: melting temperature), and the creep effect becomes important as temperature increases. A reduced strain rate introduces an increase in tensile loading time per cycle, where harmful damage is accumulated, and total test time, which results in an increase in creep damage.
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Fig. 8. The reduction of fatigue resistance with temperature and strain rate at Det ¼ 0:5%; (a) temperature dependency, (b) strain rate dependency.
When the reduction of fatigue resistance observed in the present study is explained by the effects of oxidation and creep, fatigue resistance at each temperature should reduce with decreasing strain rate, and the reduction rate in fatigue resistance with decreasing strain rate should be more severe at the higher temperatures. However, as shown in Fig. 8(b), the opposite was observed; that is, the amount of reduction in fatigue resistance with reduced strain rate decreased with an increase in temperature. This means that there is another factor, which affects fatigue resistance besides oxidation and creep, and is the DSA effect. In the investigation on Alloy 800H [2], the effects of DSA on crack initiation life and evolution of the microcrack density with the number of cycles have been studied and compared with results in vacuum, to exclude the effect of environment on crack initiation. The results reveal that DSA leads to an overall decrease in crack
initiation life, an appreciably higher crack density, and reduction in fatigue life. In other research [3,4], it has been noticed that DSA causes a faster reduction in fatigue resistance by way of rapid crack propagation where transgranular fracture is dominant. The negative strain rate stress response developed in the DSA regime introduces a higher stress concentration at the crack tip during cyclic deformation as the strain rate decreases, causing an increase in crack growth rate. This leads to a reduction in crack propagation life. It is well known that crack propagation stage is dominant in LCF conditions of ductile materials, compared with crack initiation stage, and thus the reduced crack propagation life implies a reduction in fatigue life. Therefore, it can be inferred that the steep reduction in fatigue life with reduced strain rate, in the temperature range from v 400 to 550 C (DSA regime), compared with other temperatures, is attributed to dynamic strain aging, which accelerates the reduction of fatigue resistance. In v the non-DSA regime above 600 C, the DSA effect is weakened or eliminated, and thus a slower reduction in fatigue resistance with decreasing strain rate resulted. The evidence of DSA-induced fatigue resistance reduction is manifested in Fig. 8(a). When the reduction of fatigue resistance with increasing temperature is examined at each strain rate, a steep reduction in fatigue life is observed at the high strain rate of 1 102 =s, compared with other strain rates. It is contrary to expectation since oxidation and creep effects will become more significant at low strain rate and high temperature. From the tensile test results, the regime of DSA is a function of strain rate and temperature, and moves to a higher temperature region as strain rate increases. During LCF deformation, effect v of DSA is weakened or removed beyond 600 C at the low strain rates below 1 103 =s. However, at the high strain rate of 1 102 =s, the DSA effect is still valid v even at 650 C. Hence, a severe reduction in fatigue resistance with an increase in temperature at the high strain rate of 1 102 =s can be thought to result from the DSA effect. 3.2.4. Failure mechanism in low-cycle fatigue loading conditions Fracture surfaces of LCF failed specimens were examined using SEM to study the failure mechanism in LCF loading conditions. Crack initiation occurred on the external surface in the specimen and was enhanced in the DSA regime, causing multiple crack initiation sites on the border of the fracture surface (Fig. 9(a)). v At all test conditions (400–650 C and 1 104 1 102 =s), crack propagated inward in the transgranular mode with the formation of fatigue striations (Fig. 9(b), (d) and (e)). It has been reported in SA508 [3] that dynamic strain aging would increase the degree of
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Fig. 9. SEM micrographs of LCF failed specimens at Det ¼ 0:5%; (a) 550 C, e_ ¼ 1 102 =s, (b) higher magnification view of ‘‘A’’ in (a), (c) v 650 C, e_ ¼ 1 104 =s, (d) higher magnification view of ‘‘A’’ in (c), (e) higher magnification view of ‘‘B’’ in (c). v
inhomogeniety of deformation during LCF by solute locking of moving dislocations so that DSA enhanced the partitioning of cyclic strains into segregate regions. This localized deformation leads to multiple crack initiation sites as shown in Fig. 9(a). In the study on 316L(N) stainless steel [5], where a total strain amplitude of 0.6% was employed for LCF in air, crack
initiation is transgranular and occurs from slip bands connected to the surface. Crack propagation is transgranular at higher strain rates (> 3 104 =s) in the v temperature range of 550–600 C, but at low strain rates (< 3 104 =s), intergranular cracks are noticed v on the fracture surface. Often at 600 C at these strain rates, crack propagation is assisted by oxidation. As
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shown in Fig. 9(d) and (e), the oxidation enhanced v crack propagation is observed at 650 C and strain rate of 1 104 =s, but the crack propagation mode is transgranular with fatigue striations.
4. Conclusions Based on the results and analyses on tensile and lowcycle fatigue tests, the following conclusions are made. 1. An increase in material strength and decrease in ductility were induced by dynamic strain aging in v the temperature range from 250 to 600 C, and was consistent with the regime of dynamic strain aging manifested by a plateau in the variation of dynamic strain hardening stress with temperature and a negative strain rate sensitivity. However, it shows some difference with the regime of occurrence of the serrated flow in stress–strain curves. 2. In the regime of dynamic strain aging, the fibrous zone size reduced, compared with that in other temperatures, and it is consistent with ductility reduction in the DSA regime. 3. The negative SRS regimes (DSA regimes) between tensile and LCF loading conditions coincide with each other in the temperature range from 400 to v 600 C, and SRS values under the two loading conditions are similar. 4. Softening ratio (the amount of cyclic softening during LCF) decreased in the DSA regime, compared with that in other temperatures, and it was due to the hardening by DSA. 5. A steep reduction in fatigue resistance with decreasing strain rate in the DSA regime, compared with that in other temperatures, is attributed to the enhanced crack initiation and propagation due to the DSA effect. 6. Dynamic strain aging enhanced crack initiation, causing multiple crack initiation sites, and the crack propagation mode was transgranular with fatigue v striations at all test conditions (400–650 C and 1 104 1 102 =s). Acknowledgements This work has been supported by Computer Aided Reliability Evaluation (CARE) National Research Laboratory in Korea Advanced Institute of Science and Technology (KAIST).
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