Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint

International Journal of Fatigue xxx (2015) xxx–xxx

Contents lists available at ScienceDirect

International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue

Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint Vani Shankar ⇑, K. Mariappan, R. Sandhya, K. Laha Mechanical Metallurgy Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, Tamil Nadu, India

a r t i c l e

i n f o

Article history: Received 13 April 2015 Received in revised form 31 August 2015 Accepted 1 September 2015 Available online xxxx Keywords: Low cycle fatigue Dynamic strain aging Creep–fatigue interaction 316 L(N) stainless steel Microstructure

a b s t r a c t The aim of this work is to study the time dependent effects on the low cycle fatigue (LCF) behavior of 316 L(N) stainless steel weld joint. Influence of strain rate, temperature, strain range, hold time and hold duration on fatigue life is evaluated. Occurrence of dynamic strain aging, creep damage, overall distribution of damage across the weld joint and the role of microstructure on the failure mode and failure location of the weld joint is discussed as a function of test parameters. Ó 2015 Elsevier Ltd. All rights reserved.

1. Introduction Low cycle fatigue (LCF) is an important consideration in the design of high-temperature systems subjected to thermal transients. LCF resulting from thermal transients occurs essentially under strain controlled conditions, since the surface region is constrained by the bulk of the component. In thick components the major compressive strain is introduced by the thermal transient during start up, with additional compressive and/or tensile strains during load cycling and shut down. On-load periods at elevated temperature in between transients introduce time dependent effects such as dynamic strain aging (DSA), oxidation and creep. For accurate prediction of LCF life at elevated temperatures, understanding of rate controlling time dependent damage process that influences the cyclic deformation and fracture behavior of alloys is essential. 316L(N) austenitic stainless steel is the material chosen for the primary components in liquid metal-cooled fast reactor due to its excellent high temperature mechanical properties and compatibility with the heat transfer medium i.e., liquid sodium. Fabrication of large components involves numerous weld joints and the weld joints are considered to be weak links in any structure. Since the weld joints are microstructurally and mechanically heterogeneous, these could form potential sites of failures such as fatigue failure. In addition, this heterogeneity introduces the differences in fatigue damage evolution mechanisms, fatigue crack ⇑ Corresponding author. Tel.: +91 44 274800x21147; fax: +91 44 27480075. E-mail address: [email protected] (V. Shankar).

initiation life and crack propagation rates in three zones of the weld joint, i.e. base metal, heat affected zone (HAZ) and weld metal. Therefore, weld joints are the critical sections to be considered carefully in the design of reactor components. Most of the failures have been found to originate from the HAZ in the weld joint [1–4]. Occasionally, fracture was also observed in the weld zone [3,5]. Hence, the aim of the present work is to understand the deformation and damage in 316L(N)/316L(N) SS weld joint under low cycle fatigue and creep–fatigue interaction conditions. 2. Experimental The chemical compositions in wt.% of the 316L(N) SS base metal and that of the deposited weld metal are provided in Table 1. The base metal was solution treated at 1373 K for 1 h and then water quenched in order to have both carbon and nitrogen in solid solution. Cylindrical samples of 10 mm gauge diameter and 25 mm gauge length were used for LCF and CFI tests (Fig. 1(a)). Sections of 450  250  25 mm3, cut from the mill-annealed plate were joined along the length direction by shielded metal arc welding (SMAW) process. The pads were made using type 316L(N) SS welding electrodes. The electrodes were soaked for 1 h at 473 K before the commencement of welding. During welding, the voltage and current were maintained at around 25 V and 150 A respectively. Weld joint specimens were machined from the weld pads fabricated with a double-V configuration, with an included angle of 70°, a root face of 2 mm, and a root gap of 3.15 mm. An interpass temperature of 423 K was maintained during welding. The weld

http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003 0142-1123/Ó 2015 Elsevier Ltd. All rights reserved.

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

2

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

Table 1 Chemical composition of base metal of 316 L(N) stainless steel and electrode in wt%. Materials

C

Cr

Ni

Mo

N

Mn

S

P

Fe

316 L(N) base metal 316 L(N) electrode

0.025 0.05

17.5 18.5

12.1 11.1

2.53 2.2

0.14 0.095

1.74 1.8

.0041 0.02

.017 0.025

Balance Balance

Fig. 1. Sample geometry for LCF and CFI tests (a), weld profile and location from which LCF samples were taken (b), the locations at which microstructural studies were made (1 – weld metal; 2 – heat affected zone HAZ; 3 – unaffected base metal) using EBSD (c).

pads were examined by X–radiography for their soundness. Specimens were machined from the central sections of defect-free double-V weld pads (Fig. 1(b)). Fully reversed total axial straincontrolled LCF tests were conducted at 300 K, 823 K and 873 K employing constant strain amplitude of ±0.4%, ±0.6% and ±1% and strain rate of 3  10 3 s 1, using a servohydraulic machine, equipped with a resistance heating furnace. To study dynamic strain aging behavior, LCF tests were performed at 823 K at different strain rates such as 3  10 3 s 1, 3  10 4 s 1 and 3  10 5 s 1 employing ±0.6% total strain amplitude. Creep–fatigue interaction experiments were carried out at 873 K employing ±0.6% strain amplitude and 3  10 3 s 1 strain rate using trapezoidal waveform i.e. by employing hold periods of 1, 10 and 30 min at peak tension or compression. The fatigue tested samples were sectioned parallel to the loading direction, polished, and examined under optical and field emission gun scanning electron microscopes (FEGSEM). The general microstructure of the weld joints was revealed by electrolytically etching in a solution containing 70% HNO3. Electron back scatter diffraction (EBSD) analysis was performed using FEGSEM. Automatic AztecÒ (for acquiring the EBSD data) and HKL-Channel 5TM software were used for post processing of the data. EBSD scanning voltage was 20 kV, with a scanning step size of 2 lm. Both BCC iron and FCC iron phases were used as input for indexing the EBSD patterns. A critical misorientation angle of 5° was used for inferring the presence of grain boundary between adjacent pixel points. Eight bands were used to index an EBSD pattern. Fig. 1(c) shows three different locations marked as 1, 2 and 3 (1 – weld metal, 2 – containing partly weld metal and heat affected

zone and 3 – unaffected base metal) on the 30 min tensile hold CFI tested sample at which EBSD measurements were taken. 3. Results and discussion 3.1. Initial microstructure The initial microstructure of the base and weld metal contained in the weld joint is illustrated in Fig. 2(a). The microstructure consists of equiaxed grains with some annealing twins in the base metal. The microstructure of the weld metal consists of coarse grains of dendrite like austenite structure with ferrites having typical vermicular/lacy morphologies between the dendrites as shown in Fig. 2(a). The fine duplex austenite–ferrite microstructure of the weld metal, with its many phase boundaries, offer a good resistance to the extension of fatigue cracks by causing deflection of the crack path. A typical example is depicted in Fig. 2(b). The heat affected zone adjacent to the weld metal is coarse grained austenite. Kou[6] have discussed that since the weld and the base metals have same chemical composition and crystal structure, solidification of the weld metal occurs by arranging atoms from the liquid weld metal on the partially melted substrate grains without altering their crystallographic orientations. This is clearly demonstrated in Fig. 2(a) where one observes same grains continuing in both weld metal region and in the HAZ. There is however a clear demarcation between the weld metal and the HAZ region consisting of a layer of melted region and that denotes the fusion boundary. Since the microstructure across the weld joint is quite heterogeneous, varying mechanical properties of each microstructural zone could be expected.

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

3

Fig. 2. Optical micrograph of weld joint consisting of weld metal, base metal and few grains of heat affected zone (HAZ) near the fusion line (a) and typical example to depict deflection of the fatigue crack path due to fine duplex austenite–ferrite microstructure of the weld metal (b).

3.2. Effect of temperature and strain amplitude on the cyclic stress response curve and fatigue life A comparison of the cyclic stress response curves of base metal and weld joint which were tested at 873 K at different strain amplitudes are shown in Fig. 3(a). The overall cyclic stress response curve of weld joint is higher than the base metal. The fatigue life of the weld joint is also comparatively less than the base metal tested at same test condition. As reported by Valsan et al. [1], the higher CSR of the weld joint is because of the high strength and high dislocation density present in the weld metal region. It is known that the LCF life is generally governed by the ductility of the materials at high strain amplitudes and by strength of the material at low strain amplitude. Since the weld joint consists of weld metal, HAZ and base metal, each having different mechanical properties (microhardness profile depicts varying mechanical properties across the weld joint (Fig. 3(c))), the overall fatigue life of a weld joint is governed by contribution of each microstructural zone on the overall deformation and damage. Since the weld joint contains the weld metal of high strength and hence low ductility, one observes a higher CSR and a lower fatigue life of the weld joint as compared to that of a nearly homogeneous microstructure of base metal. It is seen that generally, both base metal and weld joint exhibit similar cyclic stress response upon fatigue cycling; an initial strain

hardening to a maximum stress followed by a nearly stable peak stress and then a rapid drop in the stress value due to formation of macrocrack (Fig. 3(a) and (b)). This is similar to that reported by Kannan et al. [7] in 316L(N) base metal. The initial cyclic hardening is attributed to the individual or combined effects of (a) mutual interaction among dislocations, (b) formation of fine precipitates (e.g. Cr2N complexes, carbides) on dislocations during testing and (c) interaction between dislocations and solute atoms. Sundaraman et al. [8] emphasized on the formation of substitutional–interstitial complexes that lead to rapid initial strain hardening due to the increased resistance to dislocations motion in the matrix. Cyclic softening 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, causing an increase in the mean free path of dislocations. Valsan et al. [1] reported that gradual cyclic softening beyond saturation period has a correlation with the evolution of dislocation substructure from a random structure to a clean cell structure associated with lower stress response during cyclic deformation. Fig. 3(d) demonstrates that fatigue life decrease with increasing temperature; at lower strain amplitudes, the effect of temperature on fatigue life is more prominent. This is due to temperature and time dependent process such as oxidation. Though the microhardness profiles taken across the weld joint tested at different

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

4

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

Peak tensile stress (MPa)

500

316L(N) SS

400

300

200 -3 -1

873K,3X10 s +0.4% (BM) +0.6% (BM) +1% (BM)

100 10

+0.4% (WJ) +0.6% (WJ) +1% (WJ)

100

1000

Number of cycles

(a)

(c)

(b)

(d)

Fig. 3. (a) Comparison of cyclic stress response of base metal (BM) and weld joint (WJ) tested at 873 K, (b) effect of temperature on cyclic stress response curves of WJ, (c) microhardness profile taken across WJ tested at various temperatures and (d) effect of temperature on fatigue life of WJ.

temperatures (such as 300 K, 823 K and 873 K) does not show much change with temperature (Fig. 3(c)), there is an overall downward shift of the CSR curve with increase in temperature (Fig. 3(a) and (b)) that indicates alteration in the kinetics of initial cyclic hardening and cyclic softening in the weld joint with increasing temperature. In the following section, influence of time dependent processes such as the dynamic strain aging behavior shall be discussed. 3.3. Effect of strain rate and dynamic strain aging behavior Apart from oxidation and creep, there are other time–temperature dependent metallurgical processes such as dynamic strain aging (DSA) occurring in materials during service at elevated temperatures. Occurrence of DSA has been observed and reported recently by the present authors [9] but shall be reiterated for emphasizing the number of factors influencing fatigue life of 316 L(N) SS weld joint. The DSA phenomenon is related to the continuous locking and unlocking of mobile dislocations either by solute atmospheres or by precipitates during plastic deformation. DSA has been reported to occur under both monotonic and cyclic loading conditions. However there are distinct differences in the DSA operating temperatures under monotonic and cyclic straining. The critical temperature for the onset of dynamic strain ageing during fatigue deformation is much lower than that in monotonic tensile deformation [10] which has been attributed to the longer arrest period of dislocations in fatigue deformation in comparison

with that in monotonic deformation. Though DSA studies have been conducted on the present material [11], it shall not be included and only DSA occurrence under cyclic loading shall be discussed here. Dynamic strain aging is manifested in LCF in the form of increase in tensile peak stress with increasing temperature, increase in half-life stress with decreasing strain rate, increased cyclic hardening rate and serrations in the plastic portions of stress–strain hysteresis loops. Though serrations were not very obvious in the hysteresis loops, other manifestation of the occurrence of DSA was observed at both 823 and 873 K and hence results from 823 K are discussed below. The overall stress response increases as the strain rate is decreased from 3  10 3 s 1 to 3  10 5 s 1 (Fig. 4(a)). Though microhardness profiles taken across the weld joint at three strain rates did not show any conducive results (Fig. 4(b)), fatigue life was found to get drastically reduced with the decrease in strain rate (Fig. 4(c)). The degree or amount of hardening also increased with decrease in strain rate (Fig. 4(d)). Amount of hardening is calculated as the difference of tensile stress amplitudes for first cycle and at mid-life. During DSA, slow moving dislocations become aged by the solute atmospheres and additional dislocations are generated to maintain the imposed deformation rate. This process causes an increase in the total dislocation density. Hence the negative strain rate dependence of cyclic stress response in the DSA range results from an increase in total dislocation density during deformation. The matrix is hardened during DSA, thereby causing

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

5

Fig. 4. Effect of strain rate effect (at constant strain amplitude (±0.6%) and temperature (823 K)) on cyclic stress response (a), microhardness profile taken across the weld joint (b), fatigue life (c) and amount of hardening due to cycling (d) [9].

an increase in flow stress needed to impose the same total strain during successive cycles. It has been reported by the present authors [9] that at 823 K and 3  10 3 s 1 strain rate, transgranular mode of failure occur but under the influence of time dependent processes such as DSA, the degree of intergranularity increases; at 3  10 4 s 1 strain rate, the fraction of intergranularity is 0.19 which increases to 0.36 at slower strain rate such as 3  10 5 s 1. In the following section, effect of another time dependent process such as creep–fatigue interaction on low cycle fatigue behavior shall be discussed. 3.4. Effect of application of hold on cyclic stress response (CSR) and fatigue life In order to simulate the actual service condition, one has to apply very long hold periods (sometimes months or at least several hours (depending upon shut-down-on load period and start-up conditions during service)). However this is not practical. Hence in order to simulate the creep–fatigue interaction condition and to understand the CFI behavior under tensile and compressive holds, the longest hold period of each cycle was up to 30 min. Under the creep–fatigue interaction (CFI) condition, the period of hold in a CFI cycle denotes the time during which creep dominates; hence creep occurs during each hold period up to total CFI life. Effect of application of hold/dwell on the cyclic stress response curves are shown in Fig. 5(a). With the increase in the duration of hold, the overall CSR curve shift upwards and irrespective of

the direction of application of hold, the fatigue life decreases for all CFI tested samples as compared to the continuous cycling test. Also with the increase in the duration of hold period, the CFI life decreases and the saturation period decreases. The saturation period denotes that the rate of dislocation generation due to cyclic straining and the rate of annihilation due to occurrence of creep are balanced and a stable microstructure is attained. During the hold period, stress relaxation takes place and the elastic strains are partially converted into plastic strains. Srinivasan et al. [12] have enlisted that the plastic strain accumulated denotes the damage inculcated in the material and manifest as creep cavities at the grain boundaries and intergranular failure in materials such as 316 L(N) SS. They also supported that the comparatively lower CSR curve under CFI is due to combination of substructural recovery and the reduction in load bearing cross sectional area of the sample due to cavitated grain boundaries. Dwell sensitivity parameter (D) is shown in Fig. 5(b) for 316 L (N) SS weld joint tested under tensile hold and compressive hold for durations up to 30 min. The parameter D is calculated as the ratio of fatigue life under application of hold to the fatigue life of the material tested under continuous cycling (i.e. without application of hold). Lower value of the ratio denotes the more damaging effect of the application of hold in that direction. It is found that 316 L(N) weld joint is tensile dwell sensitive i.e. CFI life under tensile hold is lower than under equal compressive hold (Fig. 5(b)). In a review on dwell sensitive nature of materials, Goswami [13] states that the dwell sensitive nature of different metallic materials

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

6

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

Fig. 5. Effect of application of hold on (a) overall cyclic stress response (keeping constant temperature (873 K), strain amplitude (±0.6%) and strain rate (3  10 and (b) dwell sensitivity.

3

s

1

))

Fig. 6. Photograph of failed samples ((a) and (d)) and corresponding surface replica (shown on the right of the failed samples ((b) and (e)) that were tested under continuous cycling ((a), (b) and (c)) and under 10 min TH ((d) and (e)); main crack located in base metal/HAZ and shows a transgranular mode of failure (c) under CC whereas under 10 min tensile hold, replica (e) depict location of main crack in the weld metal region.

are different, depending upon the type of material, its microstructural state and testing conditions such as temperature, strain range, strain rate, hold duration and type of waveform [13] and also the environment. Materials such as, 1Cr–Mo–V [14,15], 1.25Cr–Mo [14], stainless steels; SS 304 [16,17], SS 304L [18], SS 316 [19], SS 316L [20], SS 316L(N) [21] and D9 [22] exhibit tensile dwell sensitivity. On the other hand, Brinkman et al. [23], Teranishi and McEvily [24] and Goswami [14] reported compressive dwell sensitivity in 2.25Cr–Mo. In titanium alloys Ti–6Al–4V [25] and IMI 829 [26], and many nickel base alloys [27–29] are reported to exhibit compressive dwell sensitivity. The larger compression dwell sensitivity in various alloys systems has been attributed essentially due to various phenomena, such as development of tensile mean stress, shape and size of cavities and oxidation behavior. The tensile dwell sensitivity is essentially observed in materials exhibiting creep cavitation due to a difference in the strengths of the matrix and the grain boundaries. This aspect will

be elaborated further. It is worth mentioning here that there was no development of mean stresses; in both the directions of hold application, the absolute value of mean stress did not exceed more than 5 MPa. 3.5. Cracking and failure location In order to evaluate the failure location under various testing conditions, both surface replica and microstructural analysis on the longitudinal section of the failed weld joint samples were performed. The study showed that under continuous cycling (CC), the main crack was present in the base metal (Fig. 6(a)) and the cracking mode was mainly transgranular (Fig. 6(b)). However, under the application of hold such as 10 min TH, the main crack was present in the weld metal (Fig. 6(c)). Also, the cracking mode changed from transgranular under CC (Fig. 6(b)) into intergranular under the application of hold (Fig. 7(a)–(d)). With increase in duration of

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

7

Fig. 7. Optical micrographs taken on 10 min TH sample showing cracking (red arrow marks) in the weld metal region (a), WM, HAZ and BM regions (b) and the unaffected BM (c); interlinking of cracks in the HAZ region with the cracks in the WM for sample tested under 30 min TH condition (d) and longitudinal section of failed sample showing final failure in the weld metal region (e). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

application of hold (compare 10 min TH (Fig. 7(c)) and 30 min TH (Fig. 7(d))), more grain boundary regions showed grain boundary widening and many showed decohesion (for e.g. regions marked by red arrows in Fig. 7(a)–(d)). Kim et al. [30] substantiated that cavitation is nucleated at carbides because stress is concentrated at a carbide by fatigue and grown during the hold time. These cavitated/decohesed/widened grain boundaries often link up and form larger cracks as depicted in Fig. 7(d). Grain boundary damage is typical of a tensile dwell sensitive material such as 316L(N) stainless steel and tensile dwell sensitivity (due to creep) is observed even for the short duration of hold such as 10 min (Fig. 5(b)). Mixed mode failure (that denotes creep contribution) has been reported to occur in even 1 min tensile hold [12]. Increase in the extent of intergranularity with increase in the duration of hold and creep in the form of grain boundary cavitation and decohesion has also been reported under CFI cycling in austenitic alloys [1,9,12]. It is interesting to note that even though there were numerous cracks emanating in the three microstructural zones i.e. weld metal

(Fig. 7(a)), HAZ (Fig. 7(b)) and base metal (Fig. 7(c)), the final failure occurred in the weld metal region (Fig. 7(d) and (e)) under tensile hold. It is worthwhile to mention that under continuous cycling, the microcracks form at the austenite-d-ferrite interface and the fine duplex austenite–ferrite microstructure of the weld metal, with its many phase boundaries, offer a good resistance to the extension of fatigue cracks by causing deflection of the crack path (e.g. Fig. 2(b)). Lindblom et al. [5] suggested that crack deflection causes reduced stress intensity at the crack tip with an associated reduction in the crack propagation. However, under the application of hold, there is decohesion along the austenite grain boundaries of BM, HAZ and WM and several cracks propagate along these boundaries in the weld metal. Hence, the beneficial effects of crack deflections caused by austenite-d-ferrite interface (as observed in continuous cycling) is overshadowed by joining of cracks along the austenite grain boundaries that provide an easier path for crack propagation in the weld metal. During long durations of hold (30 min TH), prolonged time is available for creep phenomenon to be operative. Hence, more number of creep cavities nucleate,

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

8

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

Fig. 8. Partitioning of typical cyclic hold hysteresis loop [33] (a) and plot of amount of stress relaxed versus accumulated plastic strain range during tensile hold CFI tests performed at 873 K on 316 L(N) SS weld joint (b).

grow and often interlink to cause crack propagation along the grain boundaries. As shown by the red arrow marks in Fig. 7(d), that most of the grain boundaries of the BM, HAZ and WM are widened/decohesced. It is apprehended that because the weld metal has a comparatively lower ductility as compared to the BM and HAZ, it is unable to accommodate the plastic strain accumulated due to stress relaxation. Hence the interlinking/joining of the cracks in the BM, HAZ and WM cause catastrophic failure in the WM (Fig. 7(d)). As stated by Nam [31] the cavities nucleate due to mechanically generated vacancies during the tensile ramp of fatigue that grow during the tensile hold by grain boundary diffusion of vacancies. According to Nam [31], the CFI life is related with the number of nucleated cavities which in turn depends upon the amount of stress relaxed, accumulated plastic strain, etc. To further understand the contribution of each microstructural zone on the overall damage process, use of EBSD is made and the results are correlated with the stress relaxation, details of which are given in the sections below. 3.6. EBSD analysis to understand deformation and damage The damage behavior in hold-time tests is generally described by characterizing the stress relaxation as a function of time elapsed in hold-time tests. Levaillant et al. [32], Wareing [33] and Rezgui et al. [34] have reported that stress relaxation occurs very rapidly during initial stages and then the rate decreases continuously at a slower rate. Rezgui et al. [34] stated that creep damage in 316 L stainless steel was produced during periods of tensile stress relaxation leading to intergranular cracking. Rodriguez and Rao [35] established that below a certain critical rate of relaxation (Fig. 8(a)), deformation is essentially due to grain boundary sliding

leading to grain boundary damage which interacts with fatigue damage to cause premature failure. The inelastic strain rates associated with rapid relaxation period (>10 4 s 1) correspond to those which are expected to cause matrix deformation while those observed in the slow stress relaxation period (<10 5 s 1) are typical of creep deformation. Hales [36] reported that as the creep component of the cycle is increased by increasing the tensile dwell period or by reducing the strain rate in cycle, the failure mode of the material under the test also changes. Plot of the amount of stress relaxed versus accumulated plastic strain shows a close to linear relationship (Fig. 8(b)). Now, the creep strain so accumulated during the CFI cycles need to be accommodated microstructurally. Since the strengths of weld metal, heat affected zone and the base metal is not the same, the accumulated strain cannot be expected to be the same in all these regions. Band contrast is a scalar value assigned to each point depending upon the brightness level of the Kikuchi diffraction pattern generated at that point and is dependent upon dislocation/crystallographic defect density and orientation. Band contrast maps are hence grey scale maps in which grain boundaries appear dark due to low pattern quality. EBSD band contrast maps taken at three different locations of 30 min tensile hold tested weld joint sample, viz. 1- weld metal, 2- containing both weld metal and heat affected zone and 3- unaffected base metal are shown in Fig. 9 A-1, A-2 and A-3 respectively. Fig. 9 A-1 shows large and elongated grains typical of directionally solidified microstructure of the weld metal. In the heat affected zone, few layers of large equiaxed grains are observed adjacent to the elongated grains of the weld metal region, Fig. 9 A-2. The unaffected base metal region show equiaxed grains, Fig. 9 A-3. In order to analyze the strain distribution across the weld joint, local misorientation strain maps corresponding to the three locations were plotted and are shown in Fig. 9 B-1, B-2 and B-3 respectively. Colors from blue1 to red denotes minimal to maximum strain respectively. Fig. 9 (B-1, B-2 and B-3) indicates that as expected, the strain is not uniform across the weld joint. Even within each microstructural zone (weld metal, HAZ or the base metal), the strain is not even. Hence there are some grains that indicate almost nil strain (blue regions) ((e.g. grains marked A, B, C in HAZ region (Fig. 9 B-2) and grains marked as D and E in BM region (Fig. 9 B-3)). Closer examination shows that deformation tends to concentrate more near the grain boundaries (yellow and red regions in Fig. 9 B-2 and B-3). Sometimes these highly strained regions, especially near the triple points of grain boundaries, also show cavities (marked by black arrows in Fig. 9 B-3(b)). There is also obvious difference in the number of regions prone to such deformation concentration at the grain boundaries for the three regions. Hence, whereas in the weld metal region, very few regions show large deformation concentration and cavities in the boundaries, that in the HAZ and base metal, there are more number of such regions. The strain maps are therefore able to qualitatively depict the extents of creep damage in the form of strain localization and cavities at the grain boundary triple points (Fig. 9). Shi and Pluvinage [20] have reported grain boundary cavitation in the base metal of 316 L(N) SS and other austenitic steels under application of tensile hold. The amount of stress relaxation occurring in each microstructural zone such as WM, HAZ and BM can be expected to be different and hence depending upon its mechanical property, each microstructural zone will accumulate different amount of plastic strain. As mentioned previously, the weld joint has high strength and high dislocation density compared to the HAZ and BM [1] and hence it would be expected that the HAZ and the BM are more yielding and can accommodate more plastic strain as compared to the weld metal. The uneven deformation/yielding/a

1 For interpretation of color in Fig. 9, the reader is referred to the web version of this article.

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

(A- 1)

9

(A- 2)

(A- 3)

(B- 1)

(B- 3 (a))

(B- 2)

(B- 3 (b))

Fig. 9. Band contrast maps ((A-1), (A-2) and (A-3)) and strain maps ((B-1), (B-2) (B-3)) for three different locations i.e. 1 – weld metal, 2 – containing both weld metal and HAZ and 3 – base metal; inset (B-3 (b)) depict strain concentration at grain boundary regions.

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003

10

V. Shankar et al. / International Journal of Fatigue xxx (2015) xxx–xxx

ccommodation of strain at each grain can be expected to be due to local constraints from the surrounding grains and the local stresses acting at each point. The unaccomodated plastic strains will give rise to formation of new surfaces in the form of cavities and cracks [37]. This is considered to be the result of microscopically irreversible accumulation of plastic strains due to dislocations pile-ups induced in the fatigue period, the creep strain accumulated due to stress relaxation in each CFI cycle and the local stress acting at each point. 4. Conclusions The study on the time independent and time dependent properties of 316 L(N) SS weld joint has shown that: (1) The cyclic stress response of the weld joint is higher than that of the base material due to the higher strength of the weld metal. (2) Dynamic strain aging manifested in the form of negative strain rate sensitivity and plays a significant role in fatigue life reduction at 823 K. (3) Creep effects are dominant at 873 K and the weld joint exhibits tensile dwell sensitivity. (4) The failure location shifts from base material and transgranular nature in continuous cycling to intergranular failure in the weld metal region during hold time tests. (5) Microstructural analysis demonstrates the role of d-ferrite– austenite interface in causing crack deflection and delay in failure, inhomogeneous strain distribution across each microstructural zone in causing cavitation at triple points in HAZ and base metal and grain boundary decohesion and joining of the cracks to cause final failure in the low ductility region of the weld metal.

Acknowledgments The authors are thankful to Dr. T. Jayakumar, Director, Metallurgy and Materials Group and Dr. A.K. Bhaduri, Associate Director, Materials Development and Technology Group, IGCAR, Kalpakkam for their support and encouragement. References [1] Valsan M, Sundaraman D, Bhanu Sankara Rao K, Mannan SL. A comparative evaluation of LCF behavior of type 316LN base metal, 316 weld metal, and 316LN/316 weld joint. Metall Mater Trans A 1995;26A:1207–19. [2] Borla Jan Z. Parent material and weldments degradation on SASOL reduction reactors due to combined effect of thermal fatigue, vibration and hydrogen attack. Int J Press Vessels Pip 2002;79:577–84. [3] Tateishi Kazuo, Hanji Takeshi. Low cycle fatigue strength of buttwelded steel joint by means of new testing system with image technique. Int J Fatigue 2004;26:1349–56. [4] Cheng Guangxu, Kaung ZB, Lou ZW, Li Hua. Experimental investigation of fatigue behavior for welded joint with mechanical heterogeneity. Int J Press Vessels Pip 1996;67:229–42. [5] Lindblom J, Linde L, Henderson PJ, Sandstorm R. Low cycle fatigue of welded joints in an austenitic stainless steel. Proc 6th int fatigue cong, vol. 2. Oxford: Pergamon; 1996. p. 1213–8. [6] Kou S. Welding metallurgy. 2nd ed. Hoboken, New Jersey: John Wiley & Sons Inc.; 2003. [7] Kannan R, Valsan M, Srinivasan VS, Bhanu Sankara Rao K, Mannan SL. Effect of temperature on LCF behavior of 316L(N) and 304L(N) stainless steels. Trans Indian Inst Met 2005;58:317–22. [8] Sundaraman D, Shankar P, Raghunathan VS. Electron microscopic study of Cr2N formation in thermally aged 316LN Austenitic stainless steels. Metall Mater Trans A 1996;27(5):1175–86.

[9] Chandra SK, Shankar V, Mariappan K, Sandhya R, Chakraborty PC. Effect of strain rate on the low cycle fatigue Behavior of 316L(N) stainless steel weld joints. Procedia Eng 2013;55:176. [10] Tsuzaki Kaneaki, Hori Takashi, Maki Tadashi, Tamura Imao. Mater Sci Eng 1983;61:247–60. [11] Shanmugavel M, Nandagopal M, Sandhya R, Bhanu Sankara Rao K, Gnanamoorthy R. Evaluation of activation energy for dynamic strain aging process in 316L(N)/316(N) SS weld joints. Trans Indian Inst Met 2008;61:301–6 [August]. [12] Srinivasan VS, Nagesha A, Valsan M, Bhanu Sankara Rao K, Mannan SL, Sastry DH. Effect of hold-time on low cycle fatigue behaviour of nitrogen bearing 316L stainless steel. Int J Press Vessels Pip 1999;76:863–70. [13] Goswami Tarun. Low cycle fatigue—dwell effects and damage mechanisms. Int J Fatigue 1999;21:55–76. [14] Goswami Tarun. Creep fatigue. I. Compilation of data and trends in the creep– fatigue behavior of low alloy steels. High Temp Mater Proc 1995;14(1):1–20. [15] Ellison EG, Patterson AFJ. Behavior of a 1Cr–Mo–V steel subject to combinations of fatigue and creep under strain control. Proc of the inst of mech engineers, vol. 190. Mechanical Eng. Publications Ltd.; 1976. p. 956–64. [16] Conway JB, Stentz RH, Berling JT. Fatigue, Tensile and Stress Relaxation Behavior of Stainless Steels. Technical Information Center, United States Atomic Energy Commission; 1975. [17] Lee Soon-Bok, Miller Alan K. A phenomenological model for intergranular fracture by r-type and wedge – type cavitation. J Eng Mater Technol 1995;117:311–21. [18] Halford GR, Johnson JR, Brown JA. High temperature LCF of Ni-201 and 304L S. S’., NASA Conference Publication 2437. In: Richmond RJ Wu ST, editor. Advanced earth-to-orbit propulsion technology; 1986. p. 172. [19] Yamaguchi K, Kanazawa K. Effect of strain wave shape on high-temperature fatigue life of a type 316 steel and application of the strain range partitioning method. Metall Trans 1980;11A:2019–27. [20] Shi HJ, Pluvinage G. Cyclic stress-strain response during isothermal and thermomechanical fatigue. Int J Fatigue 1994;16(8):549–57. [21] Srinivasan VS, Valsan M, Bhanu Sankara Rao K, Mannan SL, Raj B. Low cycle fatigue and creep–fatigue interaction behaviour of 316L(N) stainless steel and life prediction by neural network approach. Int J Fatigue 2003;25 (12):1327–38. [22] Sandhya R, Bhanu Sankara Rao K, Mannan SL. Creep–fatigue interaction behaviour of a 15Cr–15Ni, Ti modified austenitic stainless steel as a function of Ti/C ratio and microstructure. Mater Sci Eng A 2005;392(1-2):326–34. [23] Brinkman CR, Strizak JP, Brooker MK, Jaske CEJ. Time dependent strain controlled fatigue behaviour of annealed 2 1/4Cr–1Mo steel for use in nuclear steam generator design. J Nucl Mater 1976;62:181–204. [24] Teranishi H, McEvily AJ. The effect of oxidation on hold time fatigue behavior of 2.25Cr–1Mo steel. Metall Trans A 1979;Al0(11):1806–8. [25] Drapier JM, Hirchberg MH. Characterization of low cycle high temperature fatigue by the strain range partitioning method. Technical Evaluation Report of the Specialists Meeting, AGARD AR-130; 1979. p. 12. [26] Plumbridge WJ, Stanley M. Low cycle fatigue of a titanium 829 alloy. Int J Fatigue 1986;8(4):209–16. [27] Lord DC, Coffin LF. Low cycle fatigue hold time behavior of cast rene 80. Metall Trans 1973;4:1647–54. [28] Plumbridge WJ, Ellison EG. Low cycle fatigue behavior of superalloy blade materials at elevated temperature’. Mater Sci Technol 1987;3:706–15. [29] Wells CH, Sullivan CP. Interactions between creep and low-cycle fatigue in Udimet 700 at 1400 F. Fatigue High Temp, ASTM STP 1969;459:59–74. [30] Kim Dae Whan, Chang Jong-Hwa, Ryu Woo-Seog. Evaluation of the creep– fatigue damage mechanism of type 316L and Type 316LN stainless steel. Int J Press Vessels Pip 2008;85:378–84. [31] Nam SW, Yoon YC, Choi BG, Lee JM, Hong JW. The normalized Coffin–Manson plot in terms of a new damage function based on grain boundary cavitation under creep–fatigue condition. Metall Mater Trans 1996;27A:1273–81. [32] Levaillant C, Rezgui B, Pineau A. Mechanical behavior of materials. Oxfor (England): Pergamon Press; 1977 [163]. [33] Wareing J. Creep–fatigue interaction in austenitic stainless steels. Metall Trans A 1977;8A:711–21. [34] Rezgui B, Petrequin P, Mottot M. Hold time effects on LCF properties of 316 L stainless steel at 600 °C and 650 °C. Proc 5th Int conf Fract (ICF 5), vol. 5. Pergamon Press; 1982. p. 2393–402. [35] Rodriguez P, Rao KBS. Nucleation and growth of cracks and cavities under creep–fatigue interaction. Prog Mater Sci 1993;37:403–80. [36] Hales R. A quantitative metallographic assessment of structural degradation of type 316 stainless steel during creep–fatigue. Fatigue Eng Mater Struct 1980;3 (4):339–56. [37] Shankar Vani, Mariappan K, Sandhya R, Mathew MD, Jayakumar T. Effect of application of short and long holds on fatigue life of modified 9Cr–1Mo steel weld joint. Metall Trans A 2014;45(3):1390–400.

Please cite this article in press as: Shankar V et al. Understanding low cycle fatigue and creep–fatigue interaction behavior of 316 L(N) stainless steel weld joint. Int J Fatigue (2015), http://dx.doi.org/10.1016/j.ijfatigue.2015.09.003