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Thermomechanical fatigue behavior of nitrogen enhanced 316LN stainless steel: Effect of cyclic strain
International Journal of Fatigue 103 (2017) 176–184
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Thermomechanical fatigue behavior of nitrogen enhanced 316LN stainless steel: Effect of cyclic strain G.V. Prasad Reddy ⇑, A. Nagesha, R. Kannan, R. Sandhya, K. Laha Mechanical Metallurgy Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India
a r t i c l e
i n f o
Article history: Received 15 February 2017 Received in revised form 26 April 2017 Accepted 21 May 2017 Available online 23 May 2017 Keywords: Thermomechanical fatigue Cyclic stress-strain curve 316LN stainless steel Nitrogen Strain-life design curves
a b s t r a c t Thermomechanical fatigue (TMF) behavior of nitrogen enhanced 316LN stainless steel (with 0.14 wt.% N) is investigated under in-phase (IP) and out-of-phase (OP) conditions at cyclic strain amplitudes of ±0.25 to ±0.8% and with a temperature interval of 623–873 K. The study elucidates the differences in cyclic stress-strain response, TMF lives and fracture behavior of the material under IP and OP-TMF in the light of dynamic strain aging, thermal recovery and creep. The manifestations of these factors/phenomena are found to vary along stress-strain hysteresis loops, thereby leading to differences in cyclic strain hardening exponent (n0 ) and coefficient (K0 ) between the tensile and compressive branches of hysteresis loops. Irrespective of the imposed cyclic strain, transgranular fatigue failure is observed under OP-cycling in contrast to mixed-mode fracture under IP-cycling. The IP-cycling led to considerably lower cyclic life (i.e. 34–49%) compared to those under OP-TMF, with the life reduction being more significant (48.5%) at intermediate mechanical strain amplitudes of ±0.4 and ±0.6%. A comparison of the fatigue design curves under TMF with the corresponding isothermal fatigue curves (as per the RCC-MR design code) revealed a lack of conservatism in the latter. Ó 2017 Elsevier Ltd. All rights reserved.
1. Introduction In Indian sodium cooled fast reactors (SFRs), primary-side components facing hot sodium (in the temperature range 623–873 K) are currently fabricated from 316LN austenitic stainless steel (SS) with 0.06–0.08 wt.% nitrogen. During thermal transients or reactor start-up/shut-down operations, austenitic stainless steel experiences large thermal gradients due to its low thermal conductivity and high coefficient of thermal expansion. These events being repetitive during service, components are often subjected to cyclic thermal stresses. Further, thermal stresses are superimposed on primary stresses that arise from pressure loading, self-weight etc. Thus the component is subjected to both thermal cycling and mechanical strain/stress cycling, leading to thermomechanical fatigue (TMF). From the viewpoints of economy and efficiency, there exists a strong motivation for minimising the section thickness of components which in some cases may reach up to 25 mm in SFRs. Towards this end, nitrogen-enhanced 316LN SS has been envisaged for the future fast reactors of India. Nitrogen addition to 316LN SS enhances slip planarity [1–3] which in turn has been observed to improve low cycle fatigue life [2,4–6]. Apart from high tempera-
⇑ Corresponding author. E-mail address: [email protected] (G.V. Prasad Reddy). http://dx.doi.org/10.1016/j.ijfatigue.2017.05.015 0142-1123/Ó 2017 Elsevier Ltd. All rights reserved.
ture fatigue, increase in nitrogen content is also found to enhance tensile and creep properties of 316LN SS [7]. In the recent past, the beneficial effect of nitrogen on life has been found to depend on the fatigue fracture mode; nitrogen additions up to 0.14 wt.% has been observed to enhance the fatigue lives in situations where failures are predominantly transgranular [7–8]. It is important to emphasize that the literature on TMF behavior of 316LN SS with nitrogen content of 0.14 wt.% N is very limited, though some amount of work has been carried out on the alloy with a lower nitrogen content [9–12]. Further, from the perspective of generation of strain-life fatigue design curves required for designing high temperature components against TMF failure, test data over a range of strain amplitudes is essential. Hence in the present study, 316LN SS with 0.14 wt.% N is investigated over a range of mechanical strain amplitudes from ±0.25 to ±0.8%. The study presents the comparison of cyclic deformation, fracture mode and fatigue life of the material under in-phase (IP) and out-of-phase (OP) TMF cycling. Strain-life TMF design curves are generated as per the RCC-MR code, and are compared with the corresponding isothermal fatigue design curve of 316LN SS with 0.14 wt.% N and with the current RCC-MR strain-life curve of 316LN SS; RCC-MR is a French design code used for design and construction of SFR components.
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obtained from both the tensile and compressive branches of the hysteresis loop are plotted separately for IP and OP-TMF tests. In the case of IF, the stabilized or half-life hysteresis loop is more or less symmetrical about the stress and strain axes, and hence tensile stress-strain data is used for the CSS curve. In the case of OP-TMF, the two CSSCs (in Fig. 3a) correspond to ‘peak tensile stress amplitude vs. peak tensile plastic-strain amplitude data’ and ‘peak compressive stress amplitude vs. peak compressive plastic-strain amplitude data’, designated as OPT and OPC respectively. It should be mentioned that absolute values of compressive stress/strain are used for the construction of CSSCs in order to compare with the tensile stress/strain data. The position of OPT and OPC CSS curves (or IPT and IPC in Fig. 3b.) with respect to each other is justifiable from the relationship between the temperature and strain variations shown in hysteresis loops in Fig. 4 for Demech/2 of ±0.6%. It is important to note that the difference between these two CSS curves (in both Fig. 3a and b) essentially arises from the difference in tensile and compressive stress amplitudes rather than the difference in plastic strain amplitudes. Furthermore, the variation between the OPT and OPC CSS curves along the stress axis indicates the magnitude of the mean stress for a given plastic strain, and is high for the case of IP-TMF deformed samples (as apparent from Fig. 3b). It can also be noticed from Fig. 3a and b that irrespective of the strain amplitude, TMF deformation always causes a higher stress response and a lower plastic strain in comparison to the IF deformed samples. It is important to emphasize that, a similar result has also been noticed in 316LN SS (with 0.07–0.22 wt.% N) tested under IF and TMF cycling at a fixed strain rate of 3 105 s1 and at a strain amplitude of ±0.6% [10]. The extent of plastic strain spread in these tests at each mechanical strain amplitude (from ±0.25 to 1.0%) is also indicated in Fig. 3. Another important observation from Fig. 3a and b is that the OP-CSSCs (i.e. OPT and OPC curves) show relatively low cyclic strain hardening with a nearly constant slope over a range of strain amplitudes in comparison to IP-CSSCs (i.e. IPT and IPC curves). In the case of IPCSSCs, the tendency to cyclic strain hardening, however, decreases for the tests conducted at Demech/2 above ±0.6%. Indeed, the same can be observed in CSSC of IF test with a dual-slope relationship between cyclic stress and plastic-strain with a transition from one slope to other at a plastic strain (corresponding to total strain amplitude of ±0.6%). This has been found to correlate well with the corresponding change in dislocation slip mode from planar to wavy with a mixed mode of deformation at the strain amplitude of ±0.6% [3]. The IP-CSSCs also exhibits similar tendency towards dual-slope relationship, as marked by the Regimes-I and II in Fig. 3b. The cyclic strain hardening exponents are determined by fitting the CSS curves with the following power-law relationship on a log-log scale:
2. Experimental details The chemical composition of 316LN SS with 0.14 wt.% N is given in Table 1. Rectangular blanks of dimensions 155 24 22 mm, cut from the hot rolled plates along the rolling direction, were solution annealed at 1363 K for 60 min followed by water quenching. The average grain size (measured by linear intercept method) is about 86 lm. The grains are equiaxed and free from precipitates and delta-ferrite. TMF specimen geometry has a tubular crosssection with an outer diameter of 11.4 mm, 1.5 mm wall thickness and 25 mm gauge length. IP and OP-TMF tests were conducted at the mechanical strain amplitudes (Demech/2) of ±0.25, 0.4, 0.6 and 0.8% with a temperature interval (DT) of 623–873 K. All the tests were conducted in a mechanical-strain controlled mode (R = 1) with a constant strain rate of 6.4 105 s1. The cycle number corresponding to a 20% drop in peak tensile cyclic stress is taken as the fatigue life. Fracture morphologies of the tested samples were examined by scanning electron microscopy (SEM). 3. Results and discussion 3.1. Cyclic stress-strain response The cyclic deformation behavior of the material, in response to the imposed TMF cycling conditions, is discussed in terms of cyclic stress response (CSR), hysteresis loops and cyclic stress-strain curves (CSSCs). The material exhibited significant cyclic hardening, under both IP and OP-TMF cycling for a brief initial period followed by cyclic saturation/softening till the failure, except in OP test at Demech/2 = ±0.25%, as shown in tensile CSR curves in Fig. 1a and b. In the latter case, a continuous cyclic hardening till failure is observed. An increase in the strain amplitude enhanced the tendency to cyclic softening, particularly above Demech/2 of ±0.4%. As seen, the OP-TMF resulted in a higher tensile stress response in comparison to the IP case for all the strain amplitudes (Figs. 1a and b, 2a and b). Further, OP cycling resulted in a lower plastic strain range in comparison to the IP tests, as apparent from the stress-mechanical strain hysteresis loops in Fig. 2a. It is important to note the asymmetric stress-strain response in both IP and OP hysteresis loops which is due to the varying temperature throughout the imposed strain cycle. This inherent stress-strain asymmetry resulted in compressive and tensile mean stresses in the material (Fig. 2b) under IP and OP-TMF cycling respectively. The mean stress from each test, plotted as a function of the mechanical strain amplitude (Fig. 2b), shows a monotonous increase beyond the 0.25% amplitude in both IP and OP tests. Further, the magnitudes of mean stress are observed to be higher under IP-TMF compared to OP cycling. It may be noted that the mean stress and peak tensile stress in Fig. 2b are taken from the half-life loop. The cyclic stress-strain curve (CSSC), i.e. a ‘locus of half-life plastic strain amplitude vs. half-life stress amplitude’ corresponding to both IP and OP conditions is shown in Fig. 3. Also portrayed, for comparison, is the CSS curve obtained from isothermal low cycle fatigue (IF) tests conducted on cylindrical solid samples at 873 K/3 103 s1 [3]. The temperature variation along the hysteresis loop (Fig. 4) caused the differences in magnitudes of the peak stress and the peak plastic strain in the tensile and compression portions of the hysteresis loop (Fig. 2a). Hence, the CSS curves
0 Dr Dep n ¼ K0 2 2
ð1Þ
where Dr/2 is the half-life tensile stress amplitude, Dep/2 is the half-life plastic strain amplitude, n0 is the cyclic strain hardening exponent and K0 is the cyclic strength coefficient. Strain hardening exponents for IP-CSSCs are evaluated separately for the regimes above and below the stress-strain coordinate corresponding to Demech/2 of ±0.6%. In the case of OP-CSSCs, the data points satisfy a single power-law relationship. The evaluated constants are given in Table 2. As can be seen, the CSS curve under IF cycling is charac-
Table 1 The chemical composition of the nitrogen enhanced 316LN SS. Element
C
Cr
Ni
Mo
N
Mn
Si
S
P
Fe
wt.%
0.03
17.5
12.1
2.53
0.14
1.74
0.20
0.0041
0.017
Bal.
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Fig. 1. Cyclic tensile stress response curves of the material at various strain amplitudes under (a) OP-TMF and (b) IP-TMF cycling.
Fig. 2. (a) Stress-strain TMF hysteresis loops of the material tested at mechanical strain amplitudes of ±0.25%, ±0.6%, and ±0.8%, and (b) Variation of mean stress and half-life peak tensile stress as a function of mechanical strain amplitude.
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Fig. 3. Cyclic stress strain curves of the material subjected to (a) OP and (b) IP-TMF cycling, in comparison with IF data. For OPT/IPT curves the stress-strain coordinates are taken from peak tensile stress and peak tensile plastic strain amplitude, and similarly OPC/IPC are obtained from the peak compressive stress and compressive plastic strain amplitude; corresponding mechanical strain amplitudes (in%) are indicated above the x-axis line.
Fig. 4. Temperature-stress-strain distribution along the (a) OP-TMF hysteresis loop and (b) IP-TMF hysteresis loop for the tests at Demech/2 of ±0.6%. Also marked in the loops are the regimes of occurrence of DSA and creep.
Table 2 Cyclic strength coefficient (K0 ) and strain hardening exponent (n0 ) of IP and OP cyclic stress-strain curves. Regime-1
Isothermal Fatigue IPT IPC OPT OPC
Regime-2
K0
n0
K0
n0
4872 3169 2672 1658 1520
0.467 0.376 0.321 0.265 0.252
1157 723 1019 – –
0.208 0.11 0.146 – –
terized by high cyclic strength coefficient and strain hardening exponent in comparison to TMF, in spite of the higher CSR in the latter. These coefficients follow the sequence: IF > IP > OP. Further in the case of IP-CSSCs, IPT has shown higher values of K0 and n0 in Regime-I and lower values in Regime-II, in comparison to IPC. This in turn signifies the differences in contribution to cyclic deformation in compressive and tensile plastic regimes of the hysteresis loops, and is discussed in section 3.2. On the other hand in OPTMF, K0 and n0 values are more or less similar for OPC and OPTCSS curves.
3.2. Factors/phenomena influencing the cyclic stress-strain response The above differences in stress-strain response between IP and OP cycling stem from the competition between the factors/phenomena that influence the cyclic deformation over the varying temperature range in the hysteresis loops. Based on the earlier high temperature fatigue studies [3–5,10], these are identified as dynamic strain aging (DSA i.e., dislocation-solute atom interactions), thermal recovery (i.e. thermally activated dislocation motion by cross-slip and climb) and creep, apart from dislocation-dislocation interactions. Among them, DSA and dislocation-dislocation interactions contribute to cyclic strengthening while thermal recovery and creep cause cyclic strength reduction as reported in the previous studies on low cycle fatigue and creep-fatigue interaction of 316LN SS with 0.14 wt.% N [3,13– 14]. The temperature interval employed in the present TMF tests encompasses the regime of occurrence of the above factors/phenomena. The dominant one among them controls the resultant stress-strain response under TMF, thereby influencing the CSSCs. The probable regimes of predominance of DSA and creep are marked on the IP and OP hysteresis loops in Fig. 4, for Demech/2
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of ±0.6%, based on the earlier fatigue studies on the same steel [3,13–14]. In austenitic stainless steels, the temperature and strain rate regimes of occurrence of DSA correspond to 473–923 K and 105–102 s1 respectively [3,13–15]. Dynamic strain aging restricts the cross-slip and promotes dislocation motion by glide, thereby enhancing the deformation resistance [16–17]. Creep being a temperature and time-dependent process, takes place irrespective of the elastic or elastic-plastic regimes of the hysteresis loop provided the temperatures are sufficiently high and strain rates are low of the order of 105 s1 or lower. Thermal recovery and creep in type 316 SS becomes predominant at temperatures above 815 K [3,14,18]. The predominance of thermal recovery and creep usually results in attainment of maximum stress in IP/ OP hysteresis loops before the imposed mechanical strain limits (in tension/compression) are reached [19]. However, this has not been observed in the present study and the deformation resistance continuously increased up to the peak strain indicating that either DSA and/or dislocation-dislocation interactions control the overall cyclic deformation. An earlier investigation on this stainless steel established that the occurrence of DSA causes a more pronounced cyclic hardening in the temperature range 773–873 K in comparison to lower temperatures [20]. The stress response is reported to further increase with decreasing strain rate from 3 103 s1 to 3 105 s1 in the temperature 773–873 K [14]. From the temperature distribution in the tensile branch of OP hysteresis loop (Fig. 4a), it is apparent that most of the cyclic plastic deformation occurs at temperatures below 773 K where the strengthening effects associated with DSA are marginal. In fact, the magnitudes of the DSAinduced stress drops in serrations are very small or absent on the tensile/compressive branch of OP loops (Fig. 5), except on the ‘high temperature compression plastic-regime’ of 0.6% OP-loop. Hence, the tensile stress-strain response in the OPTCSSC is essentially governed by hardening due to the dislocationdislocation interactions. In order to compare the relative hardening response from the tensile and compressive plastic deformation regimes of hysteresis loop, the ‘degree of hardening’ incurred in the respective regimes is plotted separately for all the strain amplitudes (Fig. 6). The ‘degree of hardening’ is calculated over the first 100 cycles as [(r100 r1)/r1] where r represents the peak tensile or compressive stress for the cycles 1 and 100. As apparent from Fig. 6, the material showed more or less similar degree of hardening in the tensile and compressive directions under OP-TMF cycling conditions. This can also be seen from the similitude in OPC and
Fig. 6. Degree of hardening incurred by the material in the first 100 cycles, at various strain amplitudes. Degree of hardening is based on the peak tensile and compressive stresses from the IP and OP hysteresis loops.
OPT-CSS curve profiles (Fig. 3a) and nearly identical values of K0 and n0 (Table 2) extracted from the tensile and compressive branches of the OP hysteresis loops. This indicates that the stress-strain response in the entire OP loop is essentially governed by dislocation-dislocation interactions. Though DSA occurs, its effects are possibly outweighed by the hardening due to the dislocation-dislocation interactions. The role of thermal recovery and creep in ‘high temperature compressive branch of OP loop’ is to reduce the amount of hardening resulting from the dislocation-dislocation interactions, as apparent from the position of OPC with respect to the OPT curve. In the case of IP-TMF loops, tensile plastic deformation occurs at high temperatures where DSA strengthening effects are predominant and as a consequence a high degree of hardening (Fig. 6) is observed in the tensile branch of the IP-loops. Even though the low temperature compressive deformation in IP-loop raises the cyclic strength of the material (as evident from tensile and compressive stress-strain response in loops, Fig. 2a), the degree of hardening induced in the compressive branch of the loop (Fig. 6) is comparatively less than that in the tensile branch. However, the rise in cyclic strength of the material in compressive branch of the IP-loop (see IPC curve in Fig. 3b) also includes the strength contribution from DSA resulting from the previous deformation history (i.e., in the tensile portion of the loop), apart from the
Fig. 5. Serrations in TMF hysteresis loops of the material subjected to IP and OP-TMF cycling at various strain amplitudes, (a) tensile and (b) compressive portions of the loops.
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dislocation-dislocation interactions. It is important to note the similarity in profile of CSSCs of IPC and IPT, signifying that DSA controls the overall cyclic deformation in the IP loops. This can also be evidenced by the presence of serrations in both tensile and compressive plastic regimes of the IP hysteresis loops, unlike in most of the OP-loops (Figs. 5a and b). However, at the higher strain amplitude of ±0.8%, IPT and IPC CSSCs showed a reduced tendency to hardening because high cyclic strains generate significant mobile dislocation density that accommodates the imposed cyclic strain at high stresses, partially overriding the hardening resulting from DSA. Hence, K0 and n0 values in Regime-2 are comparatively lower than those in Regime-1 (Table 2). Further, in Regime-2, the values of the above coefficients are lower for the IPT compared to IPCCSSC as shown in Table 2. This is attributed to the enhanced dislocation-dislocation interactions (at high strain amplitude) in the ‘low temperature compression plastic regime’ of the stressstrain loop. 3.3. Fatigue life and fracture behavior The aforementioned differences in TMF deformation under IP and OP cycling lead to considerable variation in lives between the above two loading conditions. Fig. 7 portrays the number of cycles to failure as a function of Demech/2. As seen from the figure, cyclic lives under IP-TMF are considerably lower i.e. 34–49% compared with those of OP-TMF at all the strain amplitudes, signifying
Fig. 7. The effect of IP and OP cycling on the TMF lives of the material as a function of mechanical strain amplitude. The% reduction in IP-TMF lives with respect to the OP lives is marked on IP-TMF data.
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the deleterious influence of IP-cycling on 316LN SS. An earlier study [10] conducted at Demech/2 of ± 0.6% with a temperature interval of 573–873 K and at a low strain rate of 3 105 s1 also revealed the same (i.e. about 42% reduction in life under IP-TMF). In spite of the beneficial compressive mean stress and low tensile stresses associated with the IP-TMF deformation (Fig. 2b) compared to those developed in OP-TMF, the former proved to be detrimental in the case of 316LN stainless steels. In the present study, the deleterious effect of IP-cycling is observed to correlate well with the intergranular damage component on the fractured surfaces of the tested samples, as discussed below. The mode of fracture under OP-TMF is observed to be largely transgranular, characterized by striations. Fig. 8a and b presents the fractographs pertaining to OP-TMF tested samples at two strain amplitudes. In contrast to this, under IP-TMF cycling, significant intergranular damage in the form of decohesions of grain boundaries and triple point cracks have been noticed in addition to the transgranular striations (Fig. 9a–d), resulting in a mixed-mode failure. This indirectly implies that the presence of intergranular damage on fracture surface accelerates the propagation of overall damage that in turn causes premature failure (as evidenced by the reduced lives in IP-TMF, Fig. 7). The influence of applied cyclic strain on the IP-TMF lives depend on the extent of beneficial transgranular damage (characterized by striations) and detrimental intergranular damage (characterized by grain boundary cracks) components on the fracture surface, and hence on the interaction between them. For example, the latter can be evidenced from Fig. 10a that shows an interaction between the ‘striated primary fatigue crack’ with the surface/subsurface intergranular damage around the grains. In the present study, the extent of striated fracture surface increased with an increase in the strain amplitude (Fig. 9a–d). On the other hand, the extent of intergranular damage on the fracture surface increased with decrease in strain amplitude. Usually, cycling at low strain amplitudes (Demech/2 0.3%) typically consume large number of cycles and therefore take long time to cause significant fatigue damage (i.e. cracks and striations) during which creep deformation assumes significance in inducing the intergranular damage at most of the grain boundaries. However, the linking/coalescence of such isolated intergranular damage and its propagation would require high stress/strain amplitudes. Hence, with increase in strain amplitude above ±0.25%, the decohesion distance across the grain boundaries and the density of ‘interconnected crack networks along the grain boundaries’ increased (Fig. 10b–d). However, at high strain amplitudes, grains deform and develop transgranular damage at a rapid rate in comparison to the intergranular creep damage due to the shortage of time, and during continued cycling the former manifests in the form of striations on the fracture surfaces (Fig. 9b–d). This mitigates the
Fig. 8. SEM fractographs of the OP-TMF tested samples, illustrating characteristic striations on the fracture surface, at mechanical strain amplitudes of (a) ±0.4% and (b) ±0.8%.
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Fig. 9. SEM fractographs of the IP-TMF tested samples, illustrating mixed mode fracture characterized by intergranular damage (i.e. grain boundary cracks) and transgranular striations on the fracture surface, at mechanical strain amplitudes of (a) ±0.25%, (b) ±0.4%, (c) ±0.6% and (d) ±0.8%. Encircled regimes indicate fatigue damage (striations) and arrows represent triple point cracks and grain boundary decohesions.
Fig. 10. SEM fractographs of the IP-TMF tested samples, illustrating (a) the interaction between the propagation of surface intergranular crack (marked by arrows) and transgranular crack propagation (encircled zones), (b–d) increase in density of ‘intergranular crack networks’ with increasing strain amplitude: (b) ±0.25%, (c) ±0.4% and (d) ±0.8%.
spread of residual intergranular damage during continued cycling. Hence, due to the lack of driving force for the propagation of inter-
granular damage at low (±0.25%) and high (±0.8%) strain amplitude, the% reduction in IP-TMF lives (with respect to the OP-
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TMF) are 34% and 40% respectively in comparison to 48.5% at intermediate strain amplitudes of ±0.4 and ±0.6%. Thus, intermediate strain amplitudes cause detrimental coalescence of transgranular and intergranular damage in IP-TMF to a greater extent. The damaging effect of IP-TMF cycling arises from the in-phase relationship between the thermal and the mechanical strain cycle, particularly in the regime of the tensile plastic deformation coinciding with increasing temperature as shown in hysteresis loops (Fig. 4b). This in turn accelerates the grain boundary damage as the temperatures above 815 K (in Fig. 4b) fall within the domain of occurrence of creep deformation and DSA. It is important to mention here that, the synergistic influence of tensile creep and DSA leads to substantial intergranular damage in 316LN SS with nitrogen contents above 0.07 wt.% N [8,10,13]. DSA strengthens the grain that hardly relaxes the stress concentration near the grain boundary triple junctions, while the tensile creep weakens the grain boundaries. This causes the differences in hardening, strain accommodation and their rates both within the grain interiors and along the grain boundaries, eventually leading to strain incompatibility across the grain boundaries and triple point cracks nucleation. Indeed, such cracks have been reported in creep studies of 316 SS as a consequence of matrix hardening due to strain aging [21]. In an earlier study on 316LN SS with 0.14 wt.% N, the onset of intergranular cracking in low cycle fatigue deformation is noticed at the temperature and strain rate combination of 823 K/3 105 s1 [8]. The presence of numerous triple point cracks (i.e. wedge cracks) along the grain boundaries indicates the inherent weakness of grain boundaries resulting in interconnected intergranular cracks and grain boundary decohesion (Figs. 9 and 10). This would occur by grain boundary sliding, as most of the observed cracks are wedge cracks at the grain boundary triple junctions as shown in Figs. 9 and 10. Indeed, high resolution SEM investigation on the fracture surfaces revealed the occurrence of grain boundary sliding displacements and its link to the formation of triple point cracks and grain boundary decohesion in low strain-rate IF studies on the same stainless steel [7]. The formation of grain boundary cracks through cavity nucleation and coalescence would be minimal as it requires much lower strain rates of the order of 107 to 1010 s1 and low stresses. In contrast, such tensile creep damage is minimal/absent in the case of OP-TMF, and the contribution of creep to cyclic damage becomes negligible as it takes place in the compressive regime of OP loops, owing to the out-of-phase relationship between the thermal and strain cycles. It must be noted that compressive creep is generally beneficial as it heals the cavities/cracks that form during the course of tensile deformation [22]. In fact, several isothermal creep-fatigue interaction studies on 304 SS reported transgranular fracture in tensile plus compressive hold time tests and intergranular (or mixed mode) fracture in tensile hold tests, supporting the above rationale [22–23]. It may be mentioned in this context that slow-fast continuous cycling tests are generally found to be more detrimental causing significant intergranular damage compared to the fast-slow tests that lead to transgranular fracture in several materials (including 304 SS [24], A286 Fe-based superalloy [24] and OFHC copper [25]).
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0.08 wt.% only, and the design curve corresponding to the maximum temperature in the present study (i.e. 873 K) is shown in Fig. 11. It should be noted that these curves are generated from the IF tests performed at the peak temperature of the thermal cycle anticipated in the actual service conditions, with an assumption that such an approach has adequate in-built conservatism with respect to design. In contrast, several materials have shown lower TMF lives in comparison to the lives obtained under IF, due to the characteristic differences in deformation behavior under TMF and IF loadings [10,27–30]. In order to examine the conservatism inherent with the IF design curves (obtained as per RCC-MR code) and also to provide TMF strain-life curves for 316LN SS with high nitrogen content of 0.14%, TMF strain-life design curves are generated and compared as described in the following text. To generate the fatigue design curve under TMF, the experimental strain-life TMF data is factored by 2 on total mechanical-strain range and 20 on the number of cycles to failure, as per RCC-MR guidelines [26]. The lower sections of these curves are then joined to obtain the design curve. The above procedure has been applied for the IP and OP strain-life data and the corresponding design curves are shown in Fig. 11. The generated TMF design curves show a cross-over with respect to the current RCC-MR design curve of 316LN SS, the cross-over being dependent on the nature of TMF loading. As apparent from Fig. 11, these curves fall below the RCC-MR design curve for strain ranges above 0.35%. It can be seen that, at low strain ranges these curves either lie above the design curve or coincide with the same. Thus, the current RCC-MR design curve based on the IF data doesn’t provide a conservative life estimate for the nitrogen enhanced 316LN SS over the entire strain range. However, as mentioned above this RCC-MR curve corresponds to 316LN SS with 0.06–0.08 wt.% N. To overcome this discrepancy due to the nitrogen content, IF design curve for 316LN SS with 0.14 wt.% N is also generated as per the RCC-MR guidelines and plotted in Fig. 11. It may be noted that, for strain ranges 0.35% which are of significance for low cycle fatigue design, the IF design curve of 0.14 wt.% N doesn’t significantly differ from the existing RCC-MR curve. Here too, TMF design curves show a cross-over, and as apparent IF design curve of 0.14 wt.% N seems to be conservative only for the OP-TMF conditions that too for strain ranges below 0.5%. However, IP design curve lies far below the IF curves (including RCC-MR) indicating that the former is the most detrimental combination of temperature and strain in the case of 316LN SS, as described in the previous sections. Thus, in order to provide safe estimate of the TMF lives of the material over
3.4. Generation of strain-life design curves under TMF loading Strain-life fatigue design curves are essentially used for assessing the fatigue life of components experiencing strain-controlled cyclic deformation in service. In SFRs, 316LN SS components exposed to high temperature (>693 K) sodium and subjected to cyclic thermomechanical stresses are designed against IF failure as per the fatigue design curves established in the RCC-MR design code at various temperatures [26]. In the case of 316LN SS, the code provides fatigue design curve for the nitrogen content of 0.06–
Fig. 11. Comparison of IP and OP-TMF strain-life design curves of nitrogen enhanced 316LN SS with the corresponding isothermal fatigue (IF) design curve and RCC-MR fatigue design curve of 316LN SS (with 0.06–0.08 wt.% N).
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a large spectrum of mechanical strain ranges, it is necessary to generate TMF fatigue design curves as IF data provides a conservative life estimate only over limited strain ranges. The present study emphasizes the importance of TMF strain-life fatigue design curves towards the conservative life estimation of the nitrogen enhanced 316LN SS components subjected to TMF loading. 4. Summary In the current study, the in-phase (IP) and out-of-phase (OP) TMF behavior of nitrogen enhanced 316LN SS (0.14 wt.% N) is compared in terms of cyclic stress-strain response, fatigue life and fracture behavior, for the TMF tests conducted at cyclic strain amplitudes of ±0.25 to ±0.8% and with a temperature interval of 623–873 K. The study is summarized as follows: 1) Dynamic strain aging, thermal recovery and creep, apart from dislocation-dislocation interactions, are observed to influence the TMF deformation behavior in nitrogen enhanced 316LN SS. The manifestations of these factors/ phenomena are found to vary along the TMF hysteresis loop. 2) Cyclic strain hardening exponent (n0 ) and coefficient (K0 ), evaluated from cyclic strain-strain curves based on powerlaw relationship, differed between the tensile and compressive branches of the IP/OP hysteresis loops. 3) The mode of fracture under out-of-phase (OP) cycling is largely transgranular characterized by striations. On the other hand, in-phase (IP) cycling induced mixed-mode fracture with predominantly intergranular damage (triple point cracks and grain boundary decohesions) on the fracture surfaces. Further, an increase in strain amplitude decreased the intergranular damage and enhanced the transgranular striated failure. 4) The IP-TMF lives of the material are considerably lower (i.e. 34–49%) than those obtained under OP-TMF cycling, and the reduction in IP lives with respect to the OP is more significant (48.5%) at the intermediate mechanical strain amplitudes of ±0.4 and 0.6% compared to the two extremes of ±0.25 and ±0.8. 5) IP and OP-TMF strain-life fatigue design curves showed a cross-over with respect to the isothermal fatigue design curves (including the current RCC-MR curve), indicating that the latter provides a conservative life estimate only over limited strain ranges for the nitrogen-enhanced 316LN SS subjected to thermomechanical loading.
Acknowledgements Authors would like to thank Mr. M. Srinivasa Rao, Indira Gandhi Center for Atomic Research, Kalpakkam for the coordination in conducting the fatigue tests. References [1] Vogt JB. Fatigue properties of high nitrogen steels. J Mater Process Tech 2001;117:364–9. [2] Nilsson JO. The influence of nitrogen on high temperature low cycle fatigue behavior of austenitic stainless steels. Fatigue Fract Eng Mater Struct 1984;7:55–64. [3] Prasad Reddy GV, Sandhya R, Sankaran S, Mathew MD. Lowcycle fatigue behavior of 316LN stainless steel alloyed with varying nitrogen content: part I: cyclic deformation behavior. Metall Mater Trans A 2014;45(11):5044–56.
[4] Kim DW, Ryu WS, Hong JH, Choi SK. Effect of nitrogen on high temperature low cycle fatigue behavior in type 316L stainless steel. J Nucl Mater 1998;254:226–33. [5] Degallaix S, Degallaix G, Foct J. Influence of nitrogen solutes and precipitates on low cycle fatigue of 316L stainless steels. In: Solomon HD, Halford GR, Kaisand LR, Leis BN, editors. low cycle fatigue, ASTM STP 942, 1988, p. 798– 811. [6] Prasad Reddy GV, Sandhya R, Sankaran S, Mathew MD. Low cycle fatigue behavior of 316LN stainless steel alloyed with varying nitrogen content: Part II: fatigue life and fracture behavior. Metall Mater Trans A 2014;45 (11):5057–67. [7] Bhaduri AK, Laha K, Ganesan V, Sakthivel T, Nandagopal M, et al. Advanced materials for structural components of Indian sodium cooled fast reactors. Intl J Press Vess Piping 2016;139–140:123–36. [8] Prasad Reddy GV, Mariappan K, Kannan R, Sandhya R, Sankaran S, Bhanu Sankara Rao K. Effect of strain rate on low cycle fatigue of 316LN stainless steel with varying nitrogen content: Part-II fatigue life and fracture. Intl J Fatigue 2015;81:309–17. [9] Kim DW, Han CH, Ryu W-S. Improvement of thermomechanical fatigue life in nitrogen alloyed 316 stainless steel. Mater Sci Forum 2005;475–479:1429–32. [10] Prasad Reddy GV, Nagesha A, Sandhya R, Sankaran S, Mathew MD, Bhanu Sankara Rao K. Thermomechanical and isothermal fatigue behavior of 316LN stainless steel with varying nitrogen content. Metall Mater Trans A 2015;46:695–707. [11] Nagesha A, Valsan M, Kannan R, Bhanu Sankara Rao K, Bauer V, Christ HJ, et al. Thermomechanical fatigue evaluation and life prediction of 316L(N) stainless steel. Int J Fatigue 2009;31:636–43. [12] Nagesha A, Kannan R, Parameswaran P, Sandhya R, Bhanu Sankara Rao K, Vakil Singh. A comparative study of isothermal and thermomechanical fatigue on type 316L(N) austenitic stainless steel. Mater Sci Eng A 2010;527(21– 22):5969–75. [13] Prasad Reddy GV, Sandhya R, Sankaran S, Laha K. Creep-fatigue interaction behavior of 316LN austenitic stainless steel with varying nitrogen content. Mater Des 2015;88:972–82. [14] Prasad Reddy GV, Mariappan K, Kannan R, Sandhya R, Sankaran S, Bhanu Sankara Rao K. Effect of strain rate on low cycle fatigue of 316LN stainless steel with varying nitrogen content: Part-I Cyclic deformation behavior. Int J Fatigue 2015;81:299–308. [15] Ivanchenko Mykola. Dynamic strain aging of austenitic stainless steels and Nibase alloys. PhD thesis. Espoo, Finland: Aalto University; 2010. [16] Zhou HW, He YZ, Cui M, Cen YW, Jiang JQ. Dependence of dynamic strain ageing on strain amplitudes during the low-cycle fatigue of TP347H austenitic stainless steel at 550 °C. Int J Fatigue 2013;56:1–7. [17] Prasad K, Sarkar R, Ghosal P, Kumar V. Effects of strain waveform on low cycle fatigue behaviour of near a Timetal 834 titanium alloy. Mater Des 2011;32:1710–5. [18] Byrnes MLG, Grujicic M, Owen WS. Nitrogen strengthening of a stable austenitic stainless steel. Acta Metall 1987;35(7):1853–62. [19] Zauter R, Petry F, Christ HJ, Mugharabi H. Thermomechanical fatigue of the austenitic stainless steel AISI 304L. In: Sehitoglu H, editor. Thermomechanical fatigue behavior of materials, ASTM STP 1189, 1993, p. 70–90. [20] Prasad Reddy GV, Sandhya R, Sankaran S, Laha K. On the relationship between cyclic deformation behavior and slip mode in 316LN stainless steel with varying nitrogen content. Trans Indian Inst Met 2016;69(2):303–8. [21] Rodriguez P, Bhanu Sankara Rao K. Nucleation and growth of cracks and cavities under creep–fatigue interaction. Prog Mater Sci 1993;37:403–80. [22] Krempl E, Wundt BM. Hold time effects in high-temperature low-cycle fatigue. ASTM STP 1971;489:1–15. [23] Coffin LF. Time-dependent fatigue of structural alloys, ORNL-5073. Oak Ridge, Tenn.: Oak Ridge National Laboratory; 1977. [24] Sidey D, Coffin LF. Low cycle fatigue damage mechanisms at high temperature, in: Fong JT, editor. Fatigue mechanisms, ASTM STP 675, 1979, p. 528-68. [25] Yagi K, Tanaka C, Kubo K, Kanemar OU. Creep-fatigue interaction of SUS 316 SS under combined loading conditions. ISIJ Int 1985;25:1179–86. [26] RCC-MRx. Design and construction rules for mechanical components of nuclear installation, Section III – Tome 1 – Subsection Z: Appendix A3, AFCEN; 2012. [27] Azadi M. Effects of strain rate and mean strain on cyclic behavior of aluminum alloys under isothermal and thermo-mechanical fatigue loadings. Int J Fatigue 2013;47:148–53. [28] Ellison EG, Al-Zamily A. Stress response under thermal-mechanical strain cycling for a 1-CrMoV Ferritic steel and two 316 Stainless steels. Fatigue Fract Eng Mater Struct 1994;17(1):39–51. [29] Zamrik SY, Davis DC, Firth LC. Isothermal and thermomechanical fatigue of type 316 stainless steel, In: Verrilli MJ, Castelli MG, editors. Thermomechanical fatigue behaviour of materials, ASTM STP 1263, 1996, p. 96–116. [30] Ramesh M, Leber HJ, Janssens KGF, Diener M, Spolenak R. Thermomechanical and isothermal fatigue behavior of 347 and 316L austenitic stainless tube and pipe steels. Int J Fatigue 2011;33:683–91.
Contents lists available at ScienceDirect
International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Thermomechanical fatigue behavior of nitrogen enhanced 316LN stainless steel: Effect of cyclic strain G.V. Prasad Reddy ⇑, A. Nagesha, R. Kannan, R. Sandhya, K. Laha Mechanical Metallurgy Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India
a r t i c l e
i n f o
Article history: Received 15 February 2017 Received in revised form 26 April 2017 Accepted 21 May 2017 Available online 23 May 2017 Keywords: Thermomechanical fatigue Cyclic stress-strain curve 316LN stainless steel Nitrogen Strain-life design curves
a b s t r a c t Thermomechanical fatigue (TMF) behavior of nitrogen enhanced 316LN stainless steel (with 0.14 wt.% N) is investigated under in-phase (IP) and out-of-phase (OP) conditions at cyclic strain amplitudes of ±0.25 to ±0.8% and with a temperature interval of 623–873 K. The study elucidates the differences in cyclic stress-strain response, TMF lives and fracture behavior of the material under IP and OP-TMF in the light of dynamic strain aging, thermal recovery and creep. The manifestations of these factors/phenomena are found to vary along stress-strain hysteresis loops, thereby leading to differences in cyclic strain hardening exponent (n0 ) and coefficient (K0 ) between the tensile and compressive branches of hysteresis loops. Irrespective of the imposed cyclic strain, transgranular fatigue failure is observed under OP-cycling in contrast to mixed-mode fracture under IP-cycling. The IP-cycling led to considerably lower cyclic life (i.e. 34–49%) compared to those under OP-TMF, with the life reduction being more significant (48.5%) at intermediate mechanical strain amplitudes of ±0.4 and ±0.6%. A comparison of the fatigue design curves under TMF with the corresponding isothermal fatigue curves (as per the RCC-MR design code) revealed a lack of conservatism in the latter. Ó 2017 Elsevier Ltd. All rights reserved.
1. Introduction In Indian sodium cooled fast reactors (SFRs), primary-side components facing hot sodium (in the temperature range 623–873 K) are currently fabricated from 316LN austenitic stainless steel (SS) with 0.06–0.08 wt.% nitrogen. During thermal transients or reactor start-up/shut-down operations, austenitic stainless steel experiences large thermal gradients due to its low thermal conductivity and high coefficient of thermal expansion. These events being repetitive during service, components are often subjected to cyclic thermal stresses. Further, thermal stresses are superimposed on primary stresses that arise from pressure loading, self-weight etc. Thus the component is subjected to both thermal cycling and mechanical strain/stress cycling, leading to thermomechanical fatigue (TMF). From the viewpoints of economy and efficiency, there exists a strong motivation for minimising the section thickness of components which in some cases may reach up to 25 mm in SFRs. Towards this end, nitrogen-enhanced 316LN SS has been envisaged for the future fast reactors of India. Nitrogen addition to 316LN SS enhances slip planarity [1–3] which in turn has been observed to improve low cycle fatigue life [2,4–6]. Apart from high tempera-
⇑ Corresponding author. E-mail address: [email protected] (G.V. Prasad Reddy). http://dx.doi.org/10.1016/j.ijfatigue.2017.05.015 0142-1123/Ó 2017 Elsevier Ltd. All rights reserved.
ture fatigue, increase in nitrogen content is also found to enhance tensile and creep properties of 316LN SS [7]. In the recent past, the beneficial effect of nitrogen on life has been found to depend on the fatigue fracture mode; nitrogen additions up to 0.14 wt.% has been observed to enhance the fatigue lives in situations where failures are predominantly transgranular [7–8]. It is important to emphasize that the literature on TMF behavior of 316LN SS with nitrogen content of 0.14 wt.% N is very limited, though some amount of work has been carried out on the alloy with a lower nitrogen content [9–12]. Further, from the perspective of generation of strain-life fatigue design curves required for designing high temperature components against TMF failure, test data over a range of strain amplitudes is essential. Hence in the present study, 316LN SS with 0.14 wt.% N is investigated over a range of mechanical strain amplitudes from ±0.25 to ±0.8%. The study presents the comparison of cyclic deformation, fracture mode and fatigue life of the material under in-phase (IP) and out-of-phase (OP) TMF cycling. Strain-life TMF design curves are generated as per the RCC-MR code, and are compared with the corresponding isothermal fatigue design curve of 316LN SS with 0.14 wt.% N and with the current RCC-MR strain-life curve of 316LN SS; RCC-MR is a French design code used for design and construction of SFR components.
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obtained from both the tensile and compressive branches of the hysteresis loop are plotted separately for IP and OP-TMF tests. In the case of IF, the stabilized or half-life hysteresis loop is more or less symmetrical about the stress and strain axes, and hence tensile stress-strain data is used for the CSS curve. In the case of OP-TMF, the two CSSCs (in Fig. 3a) correspond to ‘peak tensile stress amplitude vs. peak tensile plastic-strain amplitude data’ and ‘peak compressive stress amplitude vs. peak compressive plastic-strain amplitude data’, designated as OPT and OPC respectively. It should be mentioned that absolute values of compressive stress/strain are used for the construction of CSSCs in order to compare with the tensile stress/strain data. The position of OPT and OPC CSS curves (or IPT and IPC in Fig. 3b.) with respect to each other is justifiable from the relationship between the temperature and strain variations shown in hysteresis loops in Fig. 4 for Demech/2 of ±0.6%. It is important to note that the difference between these two CSS curves (in both Fig. 3a and b) essentially arises from the difference in tensile and compressive stress amplitudes rather than the difference in plastic strain amplitudes. Furthermore, the variation between the OPT and OPC CSS curves along the stress axis indicates the magnitude of the mean stress for a given plastic strain, and is high for the case of IP-TMF deformed samples (as apparent from Fig. 3b). It can also be noticed from Fig. 3a and b that irrespective of the strain amplitude, TMF deformation always causes a higher stress response and a lower plastic strain in comparison to the IF deformed samples. It is important to emphasize that, a similar result has also been noticed in 316LN SS (with 0.07–0.22 wt.% N) tested under IF and TMF cycling at a fixed strain rate of 3 105 s1 and at a strain amplitude of ±0.6% [10]. The extent of plastic strain spread in these tests at each mechanical strain amplitude (from ±0.25 to 1.0%) is also indicated in Fig. 3. Another important observation from Fig. 3a and b is that the OP-CSSCs (i.e. OPT and OPC curves) show relatively low cyclic strain hardening with a nearly constant slope over a range of strain amplitudes in comparison to IP-CSSCs (i.e. IPT and IPC curves). In the case of IPCSSCs, the tendency to cyclic strain hardening, however, decreases for the tests conducted at Demech/2 above ±0.6%. Indeed, the same can be observed in CSSC of IF test with a dual-slope relationship between cyclic stress and plastic-strain with a transition from one slope to other at a plastic strain (corresponding to total strain amplitude of ±0.6%). This has been found to correlate well with the corresponding change in dislocation slip mode from planar to wavy with a mixed mode of deformation at the strain amplitude of ±0.6% [3]. The IP-CSSCs also exhibits similar tendency towards dual-slope relationship, as marked by the Regimes-I and II in Fig. 3b. The cyclic strain hardening exponents are determined by fitting the CSS curves with the following power-law relationship on a log-log scale:
2. Experimental details The chemical composition of 316LN SS with 0.14 wt.% N is given in Table 1. Rectangular blanks of dimensions 155 24 22 mm, cut from the hot rolled plates along the rolling direction, were solution annealed at 1363 K for 60 min followed by water quenching. The average grain size (measured by linear intercept method) is about 86 lm. The grains are equiaxed and free from precipitates and delta-ferrite. TMF specimen geometry has a tubular crosssection with an outer diameter of 11.4 mm, 1.5 mm wall thickness and 25 mm gauge length. IP and OP-TMF tests were conducted at the mechanical strain amplitudes (Demech/2) of ±0.25, 0.4, 0.6 and 0.8% with a temperature interval (DT) of 623–873 K. All the tests were conducted in a mechanical-strain controlled mode (R = 1) with a constant strain rate of 6.4 105 s1. The cycle number corresponding to a 20% drop in peak tensile cyclic stress is taken as the fatigue life. Fracture morphologies of the tested samples were examined by scanning electron microscopy (SEM). 3. Results and discussion 3.1. Cyclic stress-strain response The cyclic deformation behavior of the material, in response to the imposed TMF cycling conditions, is discussed in terms of cyclic stress response (CSR), hysteresis loops and cyclic stress-strain curves (CSSCs). The material exhibited significant cyclic hardening, under both IP and OP-TMF cycling for a brief initial period followed by cyclic saturation/softening till the failure, except in OP test at Demech/2 = ±0.25%, as shown in tensile CSR curves in Fig. 1a and b. In the latter case, a continuous cyclic hardening till failure is observed. An increase in the strain amplitude enhanced the tendency to cyclic softening, particularly above Demech/2 of ±0.4%. As seen, the OP-TMF resulted in a higher tensile stress response in comparison to the IP case for all the strain amplitudes (Figs. 1a and b, 2a and b). Further, OP cycling resulted in a lower plastic strain range in comparison to the IP tests, as apparent from the stress-mechanical strain hysteresis loops in Fig. 2a. It is important to note the asymmetric stress-strain response in both IP and OP hysteresis loops which is due to the varying temperature throughout the imposed strain cycle. This inherent stress-strain asymmetry resulted in compressive and tensile mean stresses in the material (Fig. 2b) under IP and OP-TMF cycling respectively. The mean stress from each test, plotted as a function of the mechanical strain amplitude (Fig. 2b), shows a monotonous increase beyond the 0.25% amplitude in both IP and OP tests. Further, the magnitudes of mean stress are observed to be higher under IP-TMF compared to OP cycling. It may be noted that the mean stress and peak tensile stress in Fig. 2b are taken from the half-life loop. The cyclic stress-strain curve (CSSC), i.e. a ‘locus of half-life plastic strain amplitude vs. half-life stress amplitude’ corresponding to both IP and OP conditions is shown in Fig. 3. Also portrayed, for comparison, is the CSS curve obtained from isothermal low cycle fatigue (IF) tests conducted on cylindrical solid samples at 873 K/3 103 s1 [3]. The temperature variation along the hysteresis loop (Fig. 4) caused the differences in magnitudes of the peak stress and the peak plastic strain in the tensile and compression portions of the hysteresis loop (Fig. 2a). Hence, the CSS curves
0 Dr Dep n ¼ K0 2 2
ð1Þ
where Dr/2 is the half-life tensile stress amplitude, Dep/2 is the half-life plastic strain amplitude, n0 is the cyclic strain hardening exponent and K0 is the cyclic strength coefficient. Strain hardening exponents for IP-CSSCs are evaluated separately for the regimes above and below the stress-strain coordinate corresponding to Demech/2 of ±0.6%. In the case of OP-CSSCs, the data points satisfy a single power-law relationship. The evaluated constants are given in Table 2. As can be seen, the CSS curve under IF cycling is charac-
Table 1 The chemical composition of the nitrogen enhanced 316LN SS. Element
C
Cr
Ni
Mo
N
Mn
Si
S
P
Fe
wt.%
0.03
17.5
12.1
2.53
0.14
1.74
0.20
0.0041
0.017
Bal.
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Fig. 1. Cyclic tensile stress response curves of the material at various strain amplitudes under (a) OP-TMF and (b) IP-TMF cycling.
Fig. 2. (a) Stress-strain TMF hysteresis loops of the material tested at mechanical strain amplitudes of ±0.25%, ±0.6%, and ±0.8%, and (b) Variation of mean stress and half-life peak tensile stress as a function of mechanical strain amplitude.
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Fig. 3. Cyclic stress strain curves of the material subjected to (a) OP and (b) IP-TMF cycling, in comparison with IF data. For OPT/IPT curves the stress-strain coordinates are taken from peak tensile stress and peak tensile plastic strain amplitude, and similarly OPC/IPC are obtained from the peak compressive stress and compressive plastic strain amplitude; corresponding mechanical strain amplitudes (in%) are indicated above the x-axis line.
Fig. 4. Temperature-stress-strain distribution along the (a) OP-TMF hysteresis loop and (b) IP-TMF hysteresis loop for the tests at Demech/2 of ±0.6%. Also marked in the loops are the regimes of occurrence of DSA and creep.
Table 2 Cyclic strength coefficient (K0 ) and strain hardening exponent (n0 ) of IP and OP cyclic stress-strain curves. Regime-1
Isothermal Fatigue IPT IPC OPT OPC
Regime-2
K0
n0
K0
n0
4872 3169 2672 1658 1520
0.467 0.376 0.321 0.265 0.252
1157 723 1019 – –
0.208 0.11 0.146 – –
terized by high cyclic strength coefficient and strain hardening exponent in comparison to TMF, in spite of the higher CSR in the latter. These coefficients follow the sequence: IF > IP > OP. Further in the case of IP-CSSCs, IPT has shown higher values of K0 and n0 in Regime-I and lower values in Regime-II, in comparison to IPC. This in turn signifies the differences in contribution to cyclic deformation in compressive and tensile plastic regimes of the hysteresis loops, and is discussed in section 3.2. On the other hand in OPTMF, K0 and n0 values are more or less similar for OPC and OPTCSS curves.
3.2. Factors/phenomena influencing the cyclic stress-strain response The above differences in stress-strain response between IP and OP cycling stem from the competition between the factors/phenomena that influence the cyclic deformation over the varying temperature range in the hysteresis loops. Based on the earlier high temperature fatigue studies [3–5,10], these are identified as dynamic strain aging (DSA i.e., dislocation-solute atom interactions), thermal recovery (i.e. thermally activated dislocation motion by cross-slip and climb) and creep, apart from dislocation-dislocation interactions. Among them, DSA and dislocation-dislocation interactions contribute to cyclic strengthening while thermal recovery and creep cause cyclic strength reduction as reported in the previous studies on low cycle fatigue and creep-fatigue interaction of 316LN SS with 0.14 wt.% N [3,13– 14]. The temperature interval employed in the present TMF tests encompasses the regime of occurrence of the above factors/phenomena. The dominant one among them controls the resultant stress-strain response under TMF, thereby influencing the CSSCs. The probable regimes of predominance of DSA and creep are marked on the IP and OP hysteresis loops in Fig. 4, for Demech/2
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of ±0.6%, based on the earlier fatigue studies on the same steel [3,13–14]. In austenitic stainless steels, the temperature and strain rate regimes of occurrence of DSA correspond to 473–923 K and 105–102 s1 respectively [3,13–15]. Dynamic strain aging restricts the cross-slip and promotes dislocation motion by glide, thereby enhancing the deformation resistance [16–17]. Creep being a temperature and time-dependent process, takes place irrespective of the elastic or elastic-plastic regimes of the hysteresis loop provided the temperatures are sufficiently high and strain rates are low of the order of 105 s1 or lower. Thermal recovery and creep in type 316 SS becomes predominant at temperatures above 815 K [3,14,18]. The predominance of thermal recovery and creep usually results in attainment of maximum stress in IP/ OP hysteresis loops before the imposed mechanical strain limits (in tension/compression) are reached [19]. However, this has not been observed in the present study and the deformation resistance continuously increased up to the peak strain indicating that either DSA and/or dislocation-dislocation interactions control the overall cyclic deformation. An earlier investigation on this stainless steel established that the occurrence of DSA causes a more pronounced cyclic hardening in the temperature range 773–873 K in comparison to lower temperatures [20]. The stress response is reported to further increase with decreasing strain rate from 3 103 s1 to 3 105 s1 in the temperature 773–873 K [14]. From the temperature distribution in the tensile branch of OP hysteresis loop (Fig. 4a), it is apparent that most of the cyclic plastic deformation occurs at temperatures below 773 K where the strengthening effects associated with DSA are marginal. In fact, the magnitudes of the DSAinduced stress drops in serrations are very small or absent on the tensile/compressive branch of OP loops (Fig. 5), except on the ‘high temperature compression plastic-regime’ of 0.6% OP-loop. Hence, the tensile stress-strain response in the OPTCSSC is essentially governed by hardening due to the dislocationdislocation interactions. In order to compare the relative hardening response from the tensile and compressive plastic deformation regimes of hysteresis loop, the ‘degree of hardening’ incurred in the respective regimes is plotted separately for all the strain amplitudes (Fig. 6). The ‘degree of hardening’ is calculated over the first 100 cycles as [(r100 r1)/r1] where r represents the peak tensile or compressive stress for the cycles 1 and 100. As apparent from Fig. 6, the material showed more or less similar degree of hardening in the tensile and compressive directions under OP-TMF cycling conditions. This can also be seen from the similitude in OPC and
Fig. 6. Degree of hardening incurred by the material in the first 100 cycles, at various strain amplitudes. Degree of hardening is based on the peak tensile and compressive stresses from the IP and OP hysteresis loops.
OPT-CSS curve profiles (Fig. 3a) and nearly identical values of K0 and n0 (Table 2) extracted from the tensile and compressive branches of the OP hysteresis loops. This indicates that the stress-strain response in the entire OP loop is essentially governed by dislocation-dislocation interactions. Though DSA occurs, its effects are possibly outweighed by the hardening due to the dislocation-dislocation interactions. The role of thermal recovery and creep in ‘high temperature compressive branch of OP loop’ is to reduce the amount of hardening resulting from the dislocation-dislocation interactions, as apparent from the position of OPC with respect to the OPT curve. In the case of IP-TMF loops, tensile plastic deformation occurs at high temperatures where DSA strengthening effects are predominant and as a consequence a high degree of hardening (Fig. 6) is observed in the tensile branch of the IP-loops. Even though the low temperature compressive deformation in IP-loop raises the cyclic strength of the material (as evident from tensile and compressive stress-strain response in loops, Fig. 2a), the degree of hardening induced in the compressive branch of the loop (Fig. 6) is comparatively less than that in the tensile branch. However, the rise in cyclic strength of the material in compressive branch of the IP-loop (see IPC curve in Fig. 3b) also includes the strength contribution from DSA resulting from the previous deformation history (i.e., in the tensile portion of the loop), apart from the
Fig. 5. Serrations in TMF hysteresis loops of the material subjected to IP and OP-TMF cycling at various strain amplitudes, (a) tensile and (b) compressive portions of the loops.
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dislocation-dislocation interactions. It is important to note the similarity in profile of CSSCs of IPC and IPT, signifying that DSA controls the overall cyclic deformation in the IP loops. This can also be evidenced by the presence of serrations in both tensile and compressive plastic regimes of the IP hysteresis loops, unlike in most of the OP-loops (Figs. 5a and b). However, at the higher strain amplitude of ±0.8%, IPT and IPC CSSCs showed a reduced tendency to hardening because high cyclic strains generate significant mobile dislocation density that accommodates the imposed cyclic strain at high stresses, partially overriding the hardening resulting from DSA. Hence, K0 and n0 values in Regime-2 are comparatively lower than those in Regime-1 (Table 2). Further, in Regime-2, the values of the above coefficients are lower for the IPT compared to IPCCSSC as shown in Table 2. This is attributed to the enhanced dislocation-dislocation interactions (at high strain amplitude) in the ‘low temperature compression plastic regime’ of the stressstrain loop. 3.3. Fatigue life and fracture behavior The aforementioned differences in TMF deformation under IP and OP cycling lead to considerable variation in lives between the above two loading conditions. Fig. 7 portrays the number of cycles to failure as a function of Demech/2. As seen from the figure, cyclic lives under IP-TMF are considerably lower i.e. 34–49% compared with those of OP-TMF at all the strain amplitudes, signifying
Fig. 7. The effect of IP and OP cycling on the TMF lives of the material as a function of mechanical strain amplitude. The% reduction in IP-TMF lives with respect to the OP lives is marked on IP-TMF data.
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the deleterious influence of IP-cycling on 316LN SS. An earlier study [10] conducted at Demech/2 of ± 0.6% with a temperature interval of 573–873 K and at a low strain rate of 3 105 s1 also revealed the same (i.e. about 42% reduction in life under IP-TMF). In spite of the beneficial compressive mean stress and low tensile stresses associated with the IP-TMF deformation (Fig. 2b) compared to those developed in OP-TMF, the former proved to be detrimental in the case of 316LN stainless steels. In the present study, the deleterious effect of IP-cycling is observed to correlate well with the intergranular damage component on the fractured surfaces of the tested samples, as discussed below. The mode of fracture under OP-TMF is observed to be largely transgranular, characterized by striations. Fig. 8a and b presents the fractographs pertaining to OP-TMF tested samples at two strain amplitudes. In contrast to this, under IP-TMF cycling, significant intergranular damage in the form of decohesions of grain boundaries and triple point cracks have been noticed in addition to the transgranular striations (Fig. 9a–d), resulting in a mixed-mode failure. This indirectly implies that the presence of intergranular damage on fracture surface accelerates the propagation of overall damage that in turn causes premature failure (as evidenced by the reduced lives in IP-TMF, Fig. 7). The influence of applied cyclic strain on the IP-TMF lives depend on the extent of beneficial transgranular damage (characterized by striations) and detrimental intergranular damage (characterized by grain boundary cracks) components on the fracture surface, and hence on the interaction between them. For example, the latter can be evidenced from Fig. 10a that shows an interaction between the ‘striated primary fatigue crack’ with the surface/subsurface intergranular damage around the grains. In the present study, the extent of striated fracture surface increased with an increase in the strain amplitude (Fig. 9a–d). On the other hand, the extent of intergranular damage on the fracture surface increased with decrease in strain amplitude. Usually, cycling at low strain amplitudes (Demech/2 0.3%) typically consume large number of cycles and therefore take long time to cause significant fatigue damage (i.e. cracks and striations) during which creep deformation assumes significance in inducing the intergranular damage at most of the grain boundaries. However, the linking/coalescence of such isolated intergranular damage and its propagation would require high stress/strain amplitudes. Hence, with increase in strain amplitude above ±0.25%, the decohesion distance across the grain boundaries and the density of ‘interconnected crack networks along the grain boundaries’ increased (Fig. 10b–d). However, at high strain amplitudes, grains deform and develop transgranular damage at a rapid rate in comparison to the intergranular creep damage due to the shortage of time, and during continued cycling the former manifests in the form of striations on the fracture surfaces (Fig. 9b–d). This mitigates the
Fig. 8. SEM fractographs of the OP-TMF tested samples, illustrating characteristic striations on the fracture surface, at mechanical strain amplitudes of (a) ±0.4% and (b) ±0.8%.
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Fig. 9. SEM fractographs of the IP-TMF tested samples, illustrating mixed mode fracture characterized by intergranular damage (i.e. grain boundary cracks) and transgranular striations on the fracture surface, at mechanical strain amplitudes of (a) ±0.25%, (b) ±0.4%, (c) ±0.6% and (d) ±0.8%. Encircled regimes indicate fatigue damage (striations) and arrows represent triple point cracks and grain boundary decohesions.
Fig. 10. SEM fractographs of the IP-TMF tested samples, illustrating (a) the interaction between the propagation of surface intergranular crack (marked by arrows) and transgranular crack propagation (encircled zones), (b–d) increase in density of ‘intergranular crack networks’ with increasing strain amplitude: (b) ±0.25%, (c) ±0.4% and (d) ±0.8%.
spread of residual intergranular damage during continued cycling. Hence, due to the lack of driving force for the propagation of inter-
granular damage at low (±0.25%) and high (±0.8%) strain amplitude, the% reduction in IP-TMF lives (with respect to the OP-
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TMF) are 34% and 40% respectively in comparison to 48.5% at intermediate strain amplitudes of ±0.4 and ±0.6%. Thus, intermediate strain amplitudes cause detrimental coalescence of transgranular and intergranular damage in IP-TMF to a greater extent. The damaging effect of IP-TMF cycling arises from the in-phase relationship between the thermal and the mechanical strain cycle, particularly in the regime of the tensile plastic deformation coinciding with increasing temperature as shown in hysteresis loops (Fig. 4b). This in turn accelerates the grain boundary damage as the temperatures above 815 K (in Fig. 4b) fall within the domain of occurrence of creep deformation and DSA. It is important to mention here that, the synergistic influence of tensile creep and DSA leads to substantial intergranular damage in 316LN SS with nitrogen contents above 0.07 wt.% N [8,10,13]. DSA strengthens the grain that hardly relaxes the stress concentration near the grain boundary triple junctions, while the tensile creep weakens the grain boundaries. This causes the differences in hardening, strain accommodation and their rates both within the grain interiors and along the grain boundaries, eventually leading to strain incompatibility across the grain boundaries and triple point cracks nucleation. Indeed, such cracks have been reported in creep studies of 316 SS as a consequence of matrix hardening due to strain aging [21]. In an earlier study on 316LN SS with 0.14 wt.% N, the onset of intergranular cracking in low cycle fatigue deformation is noticed at the temperature and strain rate combination of 823 K/3 105 s1 [8]. The presence of numerous triple point cracks (i.e. wedge cracks) along the grain boundaries indicates the inherent weakness of grain boundaries resulting in interconnected intergranular cracks and grain boundary decohesion (Figs. 9 and 10). This would occur by grain boundary sliding, as most of the observed cracks are wedge cracks at the grain boundary triple junctions as shown in Figs. 9 and 10. Indeed, high resolution SEM investigation on the fracture surfaces revealed the occurrence of grain boundary sliding displacements and its link to the formation of triple point cracks and grain boundary decohesion in low strain-rate IF studies on the same stainless steel [7]. The formation of grain boundary cracks through cavity nucleation and coalescence would be minimal as it requires much lower strain rates of the order of 107 to 1010 s1 and low stresses. In contrast, such tensile creep damage is minimal/absent in the case of OP-TMF, and the contribution of creep to cyclic damage becomes negligible as it takes place in the compressive regime of OP loops, owing to the out-of-phase relationship between the thermal and strain cycles. It must be noted that compressive creep is generally beneficial as it heals the cavities/cracks that form during the course of tensile deformation [22]. In fact, several isothermal creep-fatigue interaction studies on 304 SS reported transgranular fracture in tensile plus compressive hold time tests and intergranular (or mixed mode) fracture in tensile hold tests, supporting the above rationale [22–23]. It may be mentioned in this context that slow-fast continuous cycling tests are generally found to be more detrimental causing significant intergranular damage compared to the fast-slow tests that lead to transgranular fracture in several materials (including 304 SS [24], A286 Fe-based superalloy [24] and OFHC copper [25]).
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0.08 wt.% only, and the design curve corresponding to the maximum temperature in the present study (i.e. 873 K) is shown in Fig. 11. It should be noted that these curves are generated from the IF tests performed at the peak temperature of the thermal cycle anticipated in the actual service conditions, with an assumption that such an approach has adequate in-built conservatism with respect to design. In contrast, several materials have shown lower TMF lives in comparison to the lives obtained under IF, due to the characteristic differences in deformation behavior under TMF and IF loadings [10,27–30]. In order to examine the conservatism inherent with the IF design curves (obtained as per RCC-MR code) and also to provide TMF strain-life curves for 316LN SS with high nitrogen content of 0.14%, TMF strain-life design curves are generated and compared as described in the following text. To generate the fatigue design curve under TMF, the experimental strain-life TMF data is factored by 2 on total mechanical-strain range and 20 on the number of cycles to failure, as per RCC-MR guidelines [26]. The lower sections of these curves are then joined to obtain the design curve. The above procedure has been applied for the IP and OP strain-life data and the corresponding design curves are shown in Fig. 11. The generated TMF design curves show a cross-over with respect to the current RCC-MR design curve of 316LN SS, the cross-over being dependent on the nature of TMF loading. As apparent from Fig. 11, these curves fall below the RCC-MR design curve for strain ranges above 0.35%. It can be seen that, at low strain ranges these curves either lie above the design curve or coincide with the same. Thus, the current RCC-MR design curve based on the IF data doesn’t provide a conservative life estimate for the nitrogen enhanced 316LN SS over the entire strain range. However, as mentioned above this RCC-MR curve corresponds to 316LN SS with 0.06–0.08 wt.% N. To overcome this discrepancy due to the nitrogen content, IF design curve for 316LN SS with 0.14 wt.% N is also generated as per the RCC-MR guidelines and plotted in Fig. 11. It may be noted that, for strain ranges 0.35% which are of significance for low cycle fatigue design, the IF design curve of 0.14 wt.% N doesn’t significantly differ from the existing RCC-MR curve. Here too, TMF design curves show a cross-over, and as apparent IF design curve of 0.14 wt.% N seems to be conservative only for the OP-TMF conditions that too for strain ranges below 0.5%. However, IP design curve lies far below the IF curves (including RCC-MR) indicating that the former is the most detrimental combination of temperature and strain in the case of 316LN SS, as described in the previous sections. Thus, in order to provide safe estimate of the TMF lives of the material over
3.4. Generation of strain-life design curves under TMF loading Strain-life fatigue design curves are essentially used for assessing the fatigue life of components experiencing strain-controlled cyclic deformation in service. In SFRs, 316LN SS components exposed to high temperature (>693 K) sodium and subjected to cyclic thermomechanical stresses are designed against IF failure as per the fatigue design curves established in the RCC-MR design code at various temperatures [26]. In the case of 316LN SS, the code provides fatigue design curve for the nitrogen content of 0.06–
Fig. 11. Comparison of IP and OP-TMF strain-life design curves of nitrogen enhanced 316LN SS with the corresponding isothermal fatigue (IF) design curve and RCC-MR fatigue design curve of 316LN SS (with 0.06–0.08 wt.% N).
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a large spectrum of mechanical strain ranges, it is necessary to generate TMF fatigue design curves as IF data provides a conservative life estimate only over limited strain ranges. The present study emphasizes the importance of TMF strain-life fatigue design curves towards the conservative life estimation of the nitrogen enhanced 316LN SS components subjected to TMF loading. 4. Summary In the current study, the in-phase (IP) and out-of-phase (OP) TMF behavior of nitrogen enhanced 316LN SS (0.14 wt.% N) is compared in terms of cyclic stress-strain response, fatigue life and fracture behavior, for the TMF tests conducted at cyclic strain amplitudes of ±0.25 to ±0.8% and with a temperature interval of 623–873 K. The study is summarized as follows: 1) Dynamic strain aging, thermal recovery and creep, apart from dislocation-dislocation interactions, are observed to influence the TMF deformation behavior in nitrogen enhanced 316LN SS. The manifestations of these factors/ phenomena are found to vary along the TMF hysteresis loop. 2) Cyclic strain hardening exponent (n0 ) and coefficient (K0 ), evaluated from cyclic strain-strain curves based on powerlaw relationship, differed between the tensile and compressive branches of the IP/OP hysteresis loops. 3) The mode of fracture under out-of-phase (OP) cycling is largely transgranular characterized by striations. On the other hand, in-phase (IP) cycling induced mixed-mode fracture with predominantly intergranular damage (triple point cracks and grain boundary decohesions) on the fracture surfaces. Further, an increase in strain amplitude decreased the intergranular damage and enhanced the transgranular striated failure. 4) The IP-TMF lives of the material are considerably lower (i.e. 34–49%) than those obtained under OP-TMF cycling, and the reduction in IP lives with respect to the OP is more significant (48.5%) at the intermediate mechanical strain amplitudes of ±0.4 and 0.6% compared to the two extremes of ±0.25 and ±0.8. 5) IP and OP-TMF strain-life fatigue design curves showed a cross-over with respect to the isothermal fatigue design curves (including the current RCC-MR curve), indicating that the latter provides a conservative life estimate only over limited strain ranges for the nitrogen-enhanced 316LN SS subjected to thermomechanical loading.
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