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New insights to damage initiation during creep deformation of stainless steel weld joints
Materials Science & Engineering A 703 (2017) 37–44
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New insights to damage initiation during creep deformation of stainless steel weld joints
MARK
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V.D. Vijayanand , V. Ganesan, K. Laha Creep Studies Section, Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India
A R T I C L E I N F O
A B S T R A C T
Keywords: Austenitic stainless steel Weld joints Multi-pass weld joint Creep cavitation
A systematic investigation has been carried out to asses the creep cavitation behaviour in austenitic stainless steel fusion zone by performing detailed microstructural and finite element analysis. Metallographic observations of the failed cross-weld joint specimens reveal that the cavities nucleated in the near-crown region and propagated towards the near-root region. Similar pattern of nucleation and growth of cavities was observed in weld joint tested at a different stress level. The preferential nucleation of creep cavities in the near-crown region of the weld joint can be attributed to both micro and macro inhomogeneities in the weld joint. The microinhomogeneity in the fusion zone, which could be characterized using electron backscatter diffraction studies, is a result of variations in morphology and formation of thermo-mechanically treated region across the weld pass interface. The considerable strength variation between the heat affected zone, base metal and fusion zone resulted in macro-inhomogeneity in the weld joint and introduced significant stress gradients which could be illustrated using finite element analysis. Both micro and macro-inhomogeneities have a synergistic role in determining the damage initiation location in multi-pass stainless steel weld joints subjected to creep.
1. Introduction Austenitic stainless steel is the prime structural material used for fabricating fast breeder reactor's structural components [1]. This material is a unanimous choice as it exhibits superior creep strength, adequate hot corrosion resistance and has good compatibility with liquid sodium [2]. As welding is the most widely resorted technique for fabricating huge components of the reactor, evaluating the creep strength of the weld joints is also as crucial as assessing the creep properties of the base material. Creep strength of a material is usually evaluated based on its ability to resist high temperature deformation as well as microstructural damage. Both deformation and microstructural damage influence the rupture life of the material. As components in nuclear reactors operate well within the allowable design stress limits, it is the time dependent damage in the material which is the predominant life limiting factor. Creep damage in austenitic stainless steel is usually attributed to nucleation and growth of cavities [3]. Creep cavitation reduces the effective load bearing capability of the component, which may in turn result in dimensional intolerance due to the consequent localised deformation. Elaborate assessment of the creep damage which results due to cavitation in the fusion zone, can offer new perspectives for any possible microstructural and design modifications. This in turn can help mitigate the influence of cavitation
⁎
Corresponding author. E-mail address: [email protected] (V.D. Vijayanand).
http://dx.doi.org/10.1016/j.msea.2017.07.057 Received 1 June 2017; Received in revised form 17 July 2017; Accepted 18 July 2017 Available online 19 July 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
induced damage on the failure of welded austenitic stainless steel components. In case of the austenitic stainless steel base metal, cavities during creep exposure usually initiate as a result of interactions between the grain boundary sliding phenomenon and chromium carbide precipitates [3]. For shorter durations of creep exposure when the carbides are finer in size, they pin grain boundaries effectively which obstructs sliding. Prolonged elevated temperature exposure results in significant coarsening of the carbides which render them as inefficient barriers for grain boundary sliding. This results in considerable grain boundary sliding resulting in nucleation of cavities at either triple points or at the chromium carbide precipitate interface [4]. In case of the fusion zone, regions containing delta ferrite exhibit creep cavitation [5]. A certain amount of delta ferrite is intentionally prescribed in austenitic steel fusion zone to prevent hot cracking [6]. The grain size in the fusion zone is coarser than the base metal [7] and has higher grain boundary facet length making the microstructure more susceptible to cavitation. As delta ferrite phase is a metastable phase with a relatively open body-centred cubic crystal structure when compared to the close packed face-centred cubic structure of austenite, it readily transforms into brittle intermetallic phases like sigma, chi, Laves. These transformations in addition to the precipitation of chromium carbide occur at shorter time durations when compared to the
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transformations occurring in the close-packed austenite phase. As the intermetallics and precipitates coarsen, their role in resisting grain boundary sliding diminishes gradually, resulting in nucleation of cavities at the interface [8]. Further, the presence of coarser grain size in the fusion zone results in higher facet length, which reduces the threshold stress required for cavity nucleation [9]. Thus it can be stated that in case of fusion zone the transformation of delta ferrite and in case of base metal the coarsening of carbides are the crucial phenomena which influence creep cavitation in each of these respective regions. It should be noted here that in a composite austenitic stainless steel weld joint which consists of base metal and fusion zone, the precipitation of intermetallic phases and chromium carbides from delta ferrite in the fusion zone occurs more rapidly at elevated temperature exposures when compared to the precipitation and subsequent coarsening of carbides in the austenitic base metal [10]. This results in an earlier onset of creep cavitation in the fusion zone when compared to the base metal region of the weld joint. Therefore, studies pertaining to damage evolution in the fusion zone of stainless steel are more essential for assessing the performance of the welded components. Unlike the base metal, microstructure of the fusion zone is highly heterogeneous, which makes studies on creep cavity nucleation more complicated. The primary source for the formation of heterogeneities is the laying of multiple passes and the associated thermal cycling which modifies the microstructure of the weld metal which is deposited by preceding passes. It has been reported that two distinct modifications occur in the weld metal due to the deposition of the subsequent pass. One of the modifications is associated with change of delta ferrite morphology and the other is with respect to the formation of an intrinsic thermo-mechanically strengthened region [11,12]. Studies on shielded metal arc welding (SMAW) joints have shown that deposition of subsequent pass modifies the structure of delta ferrite from vermicular to globular [11]. In order to comprehensively understand the prevalence of a thermo-mechanically region, investigations were undertaken on dual pass activated tungsten inert gas (A-TIG) welded joint. As A-TIG is an autogenous welding process, the influence of delta ferrite and its morphological changes was minimised. Using impression creep testing it was shown that the strength of the first laid pass was significantly higher than that of the pass which was laid subsequently. This study clearly established that the deposition of subsequent passes strengthen the previously deposited weld layers [13]. Cavitation in the fusion zone can be linked to the nature and distribution of the inhomogeneities. In this study, a thorough investigation is undertaken to evaluate the influence of various types of inhomogeneities on creep cavitation. Based on these results the possible reasons for creep cavitation to initiate at certain specific regions of the fusion zone have been critically examined.
Table 2 Welding parameters. Electrode size (mm) Current (A) Voltage (V) Travel speed (mm/min) Heat input (kJ/mm)
Fig. 1. Schematic of weld joint specimen.
schematic of the specimen is given in Fig. 1. Uniaxial creep tests were carried out at stress levels of 140 and 200 MPa at 923 K. Automated ball indentation (ABI) testing was carried out using a spherical indenter of 0.76 mm diameter. The indentor and was made of silicon carbide. ABI tests were performed at a temperature of 923 K and speed of the cross head was maintained at 0.008 mm s−1. Impression creep (IC) tests were carried out at 923 K using a flat bottomed cylindrical indenter of 1 mm diameter. For metallographic analysis, the weld joint specimens were polished by standard metallographic procedure up to a surface finish of 1 µm. These specimens were subsequently electrolytically etched using a solution containing 60% HNO3 and 40% H20 solution at 2 V for 30 s before observation. Micro-hardness measurements were taken using Vickers hardness tester employing a load of 200 gf. Specimens for electron back scatter diffraction (EBSD) were electro-polished before observation. EBSD data was acquired using Carl Zeiss Supra 55 FEGSEM. The FEG was subjected an acceleration voltage of 20 kV. The specimen was placed in a 70° pre-tilted holder with a working distance of 16 mm. The distribution of low and high angle grain boundaries was characterized using HKL-Channel 5 software. Finite element analysis was carried out using ABAQUS version 6.11 finite element solver. The four-node three dimensional linear tetrahedron element (C3D4) was used for meshing the geometry. The mesh size was refined until convergent values for the stress values were obtained.
3. Results and discussions 3.1. As-received microstructure The macrostructure of the weld joint is the as-welded condition is shown in Fig. 2. It can be seen from the figure that the fusion zone is composed of several layers. The fusion zone consisted of dendritic austenite grains with 3–5 vol% delta ferrite. Fig. 3(a and b) show the hardness variation across the X-Xʹ (base metal-HAZ-fusion zone-HAZbase metal) and Y-Yʹ (crown to root region of the fusion zone) lines indicated in Fig. 2. It can be seen that the hardness in the fusion zone and heat affected zone are substantially higher when compared to the base metal. The fusion zone is hardened due to the restraint imposed by the adjoining base metal and the solidification stresses which causes considerable increase in dislocation density in this dendritic microstructure [14]. It has been reported that the dislocation structure in the fusion zone of stainless steel can be compared to that of the heavily cold worked structures [15]. The HAZ in stainless steel weld joints is hardened due to the repeated thermal cycling generated during the multipass welding [16]. The increase in hardness from the crown to root region (Y-Yʹ) indicates that the root region is significantly hardened due to the deposition of subsequent passes. The dashed lines in Fig. 2
2. Experimental The chemical composition of the base material and electrode is given in Table 1. The details of the SMAW parameters are furnished in Table 2. The groove angle in the base plate was 10°. The weld joint specimen was composed of the three constituents the fusion zone in the central region, flanked by the HAZ as well as the base metal regions. A Table 1 Chemical composition range of 316LN SS base metal and 316(N) electrode in wt%.
316 LN SS base metal 316 (N) SS weld metal
C
Ni
Cr
Mo
Mn
Si
S
P
N
0.025
12.15
17.57
2.53
1.74
0.2
0.004
0.017
0.14
0.051
11.2
18.45
1.94
1.37
0.48
0.005
0.024
0.13
4 128 24 97 1.89 (each pass)
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The change in morphology was observed across all weld pass interfaces from the near-crown to the near-root regions. The transformation to globular morphology occurs due to the shape instability associated with the vermicular structure [17]. In addition to the morphological change which occurs in the globular region, deposition of subsequent passes results in the formation of a thermo-mechanically treated region. As a result, globular ferrite region exhibits better deformation resistance when compared to the vermicular region. The variations in morphology and formation of a thermo-mechanically treated region create microinhomogeneity within the fusion zone. Though the modifications in morphology were evident in both the near-crown and near-root regions, the extent of the strength mismatch across the weld pass interface could vary substantially. Strength gradients across the weld pass interface in the as-welded condition can be quantified using localised misorientation measurements employing EBSD studies. Fig. 5(a and b) shows the distribution of low angle grain boundaries (LAGB) and high angle grain boundaries (HAGB) at regions adjoining the weld pass interface taken from the near-crown and near-root regions. It can be seen from the Fig. 5c that the variation in LAGB distribution across the weld pass interface is more significant in the nearcrown region. The increase in LAGB is a signature of increased local strain gradients in the material [18]. The occurrence of less prominent gradient across the interface in the near-root pass region can be attributed to homogenisation of the dislocation substructure due to the laying of subsequent passes. Though the thermal cycle modifies the microstructure of the previously laid weld passes, it tends to accumulate more restraint in regions nearer to the root pass, when multiple weld layers are deposited subsequently. The increased restraint in the root pass results in higher hardness in this region when compared to the crown region (Fig. 3b). As the number of passes which are being laid on the previous pass reduces steadily from the root to the crown direction, this restraint reduces from the root to the crown region. As a consequence, the strength gradient across the weld pass interface is more pronounced near the crown region as it is subjected to fewer thermal cycles. Hence it can be concluded that the micro-inhomogeneity across the weld pass interface is more prominent in the crown region when compared to the root region.
Fig. 2. Macrostructure of the weld joint in the as-welded condition, the dashed line shows the region along which the creep specimen was extracted. The curved lines indicate the weld pass interfaces. The continuous lines X-X′ and Y-Y′ denote lines along which hardness measurements were taken.
Fig. 3. Hardness variation across the a) longitudinal (XX′) and b) transverse direction (YY′).
3.1.2. Macro-inhomogeneity From the hardness plot (Fig. 3a) taken across the various constituents of the weld joint, it is clear that there is a significant strength gradient in the weld joint specimen. This strength gradient results in macro inhomogeneity, generating complex stress distribution in the weld joint. The influence of fusion zone in generating stress gradients is similar to those generated due to presence of a mechanical notch. For this reason, the fusion zone is termed as a “metallurgical notch” in a weld joint specimen [19]. Miniature specimen techniques were used in this study to assess the mechanical properties of the fusion zone, HAZ and base metal independently to clearly establish the nature of stress
indicate the region from where the weld joint creep specimen was extracted. Two types of inhomogeneities were investigated within the fusion zone region bounded by the dashed lines. The region in the weld joint specimen adjacent to the crown is designated as near-crown region and the region near the root pass as near-root region in this study.
3.1.1. Micro-inhomogeneity During the deposition of the subsequent passes, the morphology of the delta ferrite changes from vermicular to globular (Fig. 4(a,b,c)).
Fig. 4. a) Microstructure across the weld pass interface, the dashed line represents the weld pass interface and the regions marked V and G indicate the vermicular and globular regions respectively b) magnified micrograph of vermicular region and c) magnified micrograph of globular region.
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Fig. 5. EBSD maps showing the distribution of grain boundaries in the a) near-crown and b) near-root region c) plot of the number fraction of boundaries in the respective regions. The low angle boundaries (2–10°, marked in red) and high angle boundaries (> 10°, marked in black). The green dashed line indicates the weld pass interface. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
distribution in the specimen. The automated ball indentation testing system was used to obtain tensile properties through a stress strain microprobe. This technique can be used to obtain the load (P)-depth of penetration (dp) plot of constituent regions in the weld joint using a spherical indentor. The equivalent true stress (σt) – plastic strain (εp) plot can be obtained using suitable conversion factors which are described below.
σt =
4P πd p2 δ
(1)
dp εp = 0.2 ⎛ ⎞ ⎝D⎠ ⎜
⎟
(2)
Here δ is the material dependent constraint factor which was taken as 3.1 in this study and D is the diameter of the indenter [20]. Fig. 6 shows the variation of yield strength and ultimate tensile strength in the fusion zone, HAZ and base metal regions of the weld joint obtained from the ABI technique tested at 923 K. It can be observed that the yield strength of the fusion zone is the highest when compared to the base metal and HAZ, however the ultimate tensile strength of this region is the lowest among the three regions. The fusion zone has considerably higher dislocation density as a result of solidification stress and constraint from the base metal. Further the dislocations can generate a tangled substructure during deformation whose motion can be more effectively blocked by interphase boundaries (delta ferrite-austenite boundary) when compared to the grain boundaries [21]. The UTS values are lower in the fusion zone due to the comparatively higher degree of localised plastic deformation of delta ferrite which has more slip systems at elevated temperatures [22]. The heat affected region has higher YS and UTS when compared to the base
Fig. 6. Variation of yield stress and ultimate tensile strength across the weld joint as measured by automated ball indentation technique.
metal due to the thermo-mechanical strengthening induced by repeated thermal cycling generated during the laying of multiple passes. Impression creep curve was used to independently evaluate the creep properties of the fusion zone, HAZ and base metal regions. The punching stress (σp) and steady state penetration rate of the indentor (vi) could be related to equivalent uniaxial stress (σu) and steady state . strain rate (ε u ) respectively, using the following equations.
σu = ασp
40
(3)
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Table 4 Creep test results. Base metal
Rupture life, hours Elongation, % Failure location
⎜
⎟
(4)
Where α and β are correlation factors dependent on the material. In the present investigation the values of α and β where taken as 0.33 and 1 respectively [23]. Fig. 7 shows the impression creep curves obtained using a punching stress of 681 MPa (uniaxial stress = 225 MPa). The curve clearly shows that both the overall depth of penetration and the steady state penetration rate was significantly lower for the fusion zone. The tangled dislocation substructure in this region results in comparatively higher resistance to creep deformation. Though the HAZ and base metal exhibited almost similar steady state penetration rates, the depths of penetration in these two regions were different. The HAZ region which was subjected to repeated thermal cycling, exhibited lower deformation rate and lower depths of penetration when compared to the unaffected base metal. The impression creep tests on each of the three constituent regions were repeated for various test levels to obtain the constants A and n of Norton's constitutive equation (ε˙ = Aσ n ) which establishes the dependence of the steady state creep rate (εε̇ ˙ ) with the applied stress (σ ). The values of A and n obtained from impression creep tests are given in Table 3. 3.2. Microstructure after creep testing A comparison between the creep test results of the weld joint and base metal is given in Table 4. Rupture life and elongation of the weld joint tested at 140 MPa were lower than that of the base metal. But at the stress level of 200 MPa, the rupture life of the weld joint was considerably higher than that of the base metal. The transformation of delta ferrite to brittle intermetallics and the associated creep cavitation in the fusion zone is the cause of pre-mature failure of the weld joint tested at 140 MPa. The failure of the weld joint tested at 200 MPa occurred at the fusion zone/HAZ interface due to strain incompatibility between these two microstructurally distinct regions. The rupture life of this weld joint was considerably short when compared to that of the weld joint tested at 140 MPa, thereby limiting the influence of creep cavitation in the fusion zone. The constraint provided by the fusion zone enhanced the rupture of the weld joint tested at 200 MPa when compared to the base metal. Though the failure location of the weld
Fusion zone −1
A(MPa h n
)
2.24 × 10 4.5
−15
HAZ 6.46 × 10 3.9
Base metal −13
200 MPa
140 MPa
175 MPa
9679
537
8060
1002
34
22.5
3.5 Within fusion zone
19 HAZ/fusion zone interface
3.2.1. Influence of micro-inhomogeneity Due to microstructural variations across the weld pass interface in the crown region, there is significant strain partitioning between the globular and vermicular delta ferrite regions under the influence of external stress. The strain partitioning due to the strength mismatch between the vermicular and globular ferrite region results in preferential nucleation of cavities in the vermicular region. The propagation of these cavities are however restricted at the weld pass interface. The globular delta ferrite region on the other side of the weld pass interface offers more resistance to propagation of cavities. The weld pass interface in the crown region obstructs the propagation of cavities. It can also be stated that in this region the interface serves as a boundary separating the cavity prone and cavity resistant regions. But in the root region the strength mismatch between the vermicular and globular delta ferrite is not significant. Therefore, under the influence of an external stress, strain partitioning is not prominent between the vermicular and globular ferrite regions in the root region. The absence of strain partitioning circumvents cavity propagation across the weld pass interfaces in the root region which renders it more resistant to creep cavitation.
Table 3 Values of Norton's constants A and n.
-n
140 MPa
joint varied with applied stress, considerable cavitation could be observed under both the test conditions. The macrostructure of the weld joint specimen tested at a stress level of 140 MPa at 923 K is shown in Fig. 8a. Fig. 8b shows a magnified region of the fusion zone near the weld pass interface. Preferential cavitation could be seen in the vermicular region of the weld pass interface. During creep exposure, the vermicular ferrite region near the weld pass interface is more susceptible to creep damage due to the favourable kinetics for transformation of delta ferrite to the brittle intermetallic phases [8]. The transformations are more sluggish in the globular morphology and this strengthened region offers more resistance to crack propagation and as a consequence, the cracks are blocked at the weld pass interface. The macrostructure also shows that in addition to the predominant crack which resulted in ultimate failure, there was another interlinked crack propagating to more than half of the specimen diameter. This suggests that the failure initiated in the near-crown region of the weld joint and propagated towards the root side. The macrostructure of the weld joint tested at higher stress level of 200 MPa is shown in Fig. 9. Since the failure of this joint occurred at the fusion zone/HAZ interface, damage distribution in the fusion zone was evident more explicitly. The density of cavities in this weld joint was quite high near the crown side and reduced steadily near the root side. Similar observation could be made for the joint tested at 140 MPa, where significant number of un-propagated cavities was observed in the crown region when compared to the root region. But there was a distinct variation in the location of the cavities with respect to distance from the weld centre line for the weld joint tested at 200 MPa. It could be clearly seen that the fusion zone cavitation in this weld joint occurred at regions adjacent to the HAZ interface, whereas for the weld joint tested at 140 MPa, the propagated cracks were almost along the weld centre line. The influence of micro and macro inhomogeneities on the cavity initiation and propagation is assessed in the following sections.
Fig. 7. Variation of depth of penetration with time for the three constituent regions as obtained from impression creep testing. . v ε u=⎛ i ⎞ βd ⎝ ⎠
Weld joint
1.05 × 10−12 3.9
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Fig. 8. a) Macrostructure of weld joint tested at 140 MPa/923 K, region A and B marked in the figure denote the crack initiation and final failure locations respectively, the arrows indicate the direction of the applied stress b) microstructure showing preferential cavitation in the same weld joint.
analysis. As Norton's law is only applicable in the secondary creep regime, the simulation was carried out for duration much shorter than the apparent secondary creep regime in the weld joint. This approach was adopted because from the creep curve of a composite weld joint it is not feasible to determine the relative contributions of the three constituent regions to the cumulative creep strain. Any possible localised microstructural or geometrical changes in any of the three regions would result in deviation from steady state behaviour. In order to offset any localised effect in the weld joint, the simulation was carried for duration of 6000 and 750 h, much lower than the rupture life of the specimen tested at 140 MPa and 200 MPa respectively. The variation of the stress tri-axiality factor obtained from the FE simulation was examined in this study due to its definitive role in influencing creep cavitation behaviour. Stress tri-axiality factor can be estimated using the following relationship.
Fig. 9. Macrostructure of weld joint tested at 200 MPa/923 K. The arrows indicate the direction of the applied stress.
3.2.2. Influence of macro-inhomogeneity The presence of microstructurally varying constituents results in significant stress gradients across the weld joint, which get amplified due to the presence of the groove angle. In order to study the evolution of the state of stress during creep, finite element analysis was performed on simulated weld joint geometry. One half of the composite weld joint specimen was modelled for the study and the simulation was performed under external stress levels of 140 and 200 MPa. The partitioning of the geometry into the fusion zone, heat affected zone and base metal was done based on the microstructure and hardness readings (Fig. 10). The four transverse lines indicated in the figure denote the paths along which the stress gradients were analysed. Path 1 was drawn along the weld centre line and paths 2, 3 and 4 were separated by a constant distance of 2 mm. Path 4 was closest to the fusion zone and HAZ interface. The tensile and creep properties used in the simulation were obtained by automated ball indentation and impression creep testing techniques. Norton's constitutive equation was used to simulate the plastic behaviour of the material under creep in the finite element
Stress Tri−axiality Factor =
σ1 + σ2 + σ3 3 × σvm
(5)
Where σ1, σ2, σ3 are the principal stress components and σvm is the vonMises stress. It has been well documented that both the von-Mises stress and the maximum principal stress influence nucleation and growth of cavities [24]. The tri-axiality factor balances the influence of both these factors. Fig. 11a and b show the distribution of stress tri-axiality factor for the weld joint geometries simulated at stress levels of 140 and 200 MPa. It could be seen from these figures that the presence of microstructurally varying constituent regions generate a significant variation in the state of stress across the weld joint. The influence of groove angle on the state of stress could be more clearly visualised in Fig. 12a and b. For the weld joint tested at both the stress levels, the tri-axial state of stress along transverse direction is almost constant near the weld centre line and starts to show a gradient with increase in distance from weld centre line. The variation in the tri-axiality factor in Paths 1, 2 and 3 was almost symmetric about the specimen's longitudinal axis. This trend could possibly be attributed to the relatively small groove angle (10°), which did not influence the variation of tri-axiality factor along these paths. Whereas for the path adjoining the interface (Path 4), the stress tri-axiality showed significant gradient with a minimum near the longitudinal axis of the specimen and increase in values along both the radial directions. It could also be seen that stress tri-axiality factor was marginally higher in the crown region when compared to the root region. Thus depending on the distance from the weld centre line, the presence of the groove angle can generate can generate stress gradients along the crown to root direction. In case of the weld joint tested at 140 MPa, the micro-inhomogeneity caused by the strength mismatch across the weld pass interface is more dominant in initiating creep cavitation in the crown region. The variation of stress tri-axiality factor resulting due to the
Fig. 10. Partitioned weld joint based on the microstructure and hardness. The red transverse lines in the figure indicate the paths along which the tri-aixiality factors are subsequently plotted. The blue longitudinal dashed line indicates the specimen's longitudinal axis. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
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Fig. 11. Contours of stress tri-axiality factor for weld joint simulated with applied stress of a) 140 MPa and b) 200 MPa.
comprising dimples which is a signature of ductile failure (Fig. 13a, b and c). The regions of the fractograph could be well correlated to the corresponding optical micrograph (Fig. 8a). Inter-dendritic failure can be linked to the propagation of cavities along the dendritic boundaries where prior transformed delta ferrite phase segregates during solidification. During creep exposure, the brittle intermetallics which transform from delta ferrite are also found segregated at the inter-dendritic boundaries. For the joint under the current investigation the cavity propagation occurred in the region marked A which corresponds to the near-crown region. With creep exposure, the propagation occurred into the specimen by possible inter-linking with other cavities nucleated in the bulk of the specimen. After significant propagation of the cracks the final failure of the specimen occurred in region B. Region B in the fractograph which showed evidences of ductile failure corresponds to the failure mode in the near-root pass region. The region adjacent to the root pass is strengthened significantly due to the laying of subsequent passes as a result it offered more resistance to inter-dendritic cavity propagation. The ductile failure in this region was a result of significant plastic deformation caused by stress ‘overload’. This is because the propagation of cavities causes significant reduction in load bearing cross sectional area and the final ductile failure is a result of significant plastic deformation, due to the associated stress concentration in the root pass region
groove angle could not have possibly influenced creep cavitation as the failure location was almost near the weld centre line. It can be seen that the distribution and values of tri-axiality factor in the fusion zone were similar for both the weld joint tested at two different stress levels. However, as the associated von-Mises stress (which is dependent on the applied stress) is much lower in case of the weld joint specimen tested at 140 MPa, the possibility of nucleation of cavities due to the geometrical constraint is inconsequential. Under the influence of a higher stress level of 200 MPa, the effects of macro-inhomogeneity are magnified causing preferred nucleation in the crown region of the fusion zone adjoining the HAZ interface. Even in the crown region, the microinhomogeneity due to the strength mismatch resulted in cavitation more specifically in the vermicular region. Therefore, in case of the weld joint tested at 200 MPa, both micro and macro-inhomogeneities played a synergistic role in initiating creep cavitation. It should also be noted in this weld joint that though the stress tri-axiality factor in the crown region was highest along path 4 (Fig. 12b), cavitation occurred only at a finite distance from the HAZ interface. This is because, only in these regions the grain boundaries oriented to almost 45°, which is the direction most favourable for stress assisted grain boundary sliding to occur [25]. It can also be inferred from this study that a sharper groove angle has the potential to cause significant variation in the stress tri-axiality factor across the transverse direction at shorter distances from the weld centre-line. This can cause significant increase in the tri-axiality factor in the crown region thereby catalysing the preferential nucleation of cavities in this region.
4. Conclusion Creep failure in stainless steel weld joints initiate in the crown region and propagate into the root region. The damage initiation in the crown region can be attributed to two inhomogeneities – a) micro-inhomogeneity caused due to the variation in strength gradient in the crown and the root region and b) macro-inhomogeneity due to the presence of a groove angle which changes fusion zone geometry from crown to root. The micro-inhomogeneity is the predominant factor causing preferential damage initiation in the weld joint at a stress level
3.3. Fractography The location of crack initiation site, crack propagation and final failure generated distinct features in the fractograph of the weld joint specimen tested at 140 MPa. The fractograph composed of a predominantly intergranular dendritic failure and a small region
Fig. 12. Variation of stress tri-aixiality factor taken along Paths 1, 2, 3, 4 for weld joint simulated with applied stress of a) 140 MPa and b) 200 MPa.
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Fig. 13. Fractograph of specimen failed at 140 MPa, region A and B marked in the figure denote the crack initiation and final failure locations respectively, b) magnified image of region A and c) magnified image of region B. [9] A.G. Evans, Acta Metall. 28 (1980) 1155–1163. [10] J.J. Smith, R.A. Farrar, Int. Mater. Rev. 38 (1993) 25–51. [11] V.D. Vijayanand, K. Laha, P. Parameswaran, V. Ganesan, M.D. Mathew, Mater. Sci. Eng. A 607 (2014) 138–144. [12] V.D. Vijayanand, J. Ganesh Kumar, P.K. Parida, V. Ganesan, K. Laha, Metall. Mater. Trans. A 48 (2017) 706–721. [13] V.D. Vijayanand, M. Vasudevan, V. Ganesan, P. Parameswaran, K. Laha, A.K. Bhaduri, Metall. Mater. Trans. A 47 (2016) 2804–2814. [14] T.G. Gooch, J. Honeycombe, Weld. J. 59 (1980) 233–241. [15] M. Valsan, D. Sundararaman, K.B.S. Rao, S.L. Mannan, Metall. Mater. Trans. A 26 (1995) 1207–1220. [16] C.F. Etienne, D.V. Rossum, F. Roode, in: Proceedings of the International Conference on Engineering Aspects of Creep, University of Sheffield, 2, 1980, pp. 113–121. [17] S.A. David, Weld. J. 60 (1981) 63–71. [18] S.I. Wright, M.M. Nowell, D.P. Field, Microsc. Microanal. 17 (2011) 316–329. [19] M.J. Manjoine, Weld. J. 61 (1982) 50–57. [20] J. Ganesh Kumar, V.D. Vijayanand, M. Nandagopal, K. Laha, Mater. High Temp. 32 (2015) 619–626. [21] E. Werner, H.P. Stuwe, Mater. Sci. Eng. A 68 (1984 & ;;85) 175–182. [22] C.R. Weinberger, B.L. Boyce, C.C. Battaile, Int. Mater. Rev. 58 (2013) 296–314. [23] Naveena, V.D. Vijayanand, V. Ganesan, K. Laha, M.D. Mathew, Mater. Sci. Eng. A 552 (2012) 112–118. [24] T.L. Sham, A. Needleman, Acta Metall. 31 (1993) 919–926. [25] R.C. Gifkins, A. Gittins, R.L. Bell, T.G. Langdon, J. Mater. Sci. 3 (1968) 306–313.
of 140 MPa. In addition to this reason, the variation in macro-inhomogeneity could also influence creep cavitation for the weld joint tested at a higher stress level of 200 MPa. Though cavitation in the fusion zone of stainless steel weld joints occurs in the crown region, depending on the applied stress and groove angle its location from the weld centre line can vary substantially. References [1] S.L. Mannan, S.C. Chetal, Baldev Raj, S.B. Bhoje, Trans. Indian Inst. Met. 235 (2005) 155–178. [2] Baldev Raj, P. Chellapandi, P.R. Vasudeva Rao, Sodium Fast Reactors with Closed Fuel Cycles, CRC Press, Taylor and Francis Group, Florida, 2015, pp. 86–90. [3] I.W. Chen, A.S. Argon, Acta Metall. 29 (1981) 1321–1333. [4] G. Sasikala, M.D. Mathew, K.B.S. Rao, S.L. Mannan, Metall. Mater. Trans. A 31 (2000) 1175–1185. [5] C.A.P. Horton, P. Marshall, R.G. Thomas, Mechanical Behavior and Nuclear Applications of Stainless Steel at Elevated Temperatures, The Metals Society, London, 1982, pp. 66–72. [6] K.H. Lo, C.H. Shek, J.K.I. Lai, Mater. Sci. Eng. R. 65 (2009) 39–104. [7] J.C. Lippold, W.F. Savage, Weld. J. 59 (1980) 48–58. [8] B.A. Senior, J. Mater. Sci. 25 (1990) 45–53.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
New insights to damage initiation during creep deformation of stainless steel weld joints
MARK
⁎
V.D. Vijayanand , V. Ganesan, K. Laha Creep Studies Section, Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India
A R T I C L E I N F O
A B S T R A C T
Keywords: Austenitic stainless steel Weld joints Multi-pass weld joint Creep cavitation
A systematic investigation has been carried out to asses the creep cavitation behaviour in austenitic stainless steel fusion zone by performing detailed microstructural and finite element analysis. Metallographic observations of the failed cross-weld joint specimens reveal that the cavities nucleated in the near-crown region and propagated towards the near-root region. Similar pattern of nucleation and growth of cavities was observed in weld joint tested at a different stress level. The preferential nucleation of creep cavities in the near-crown region of the weld joint can be attributed to both micro and macro inhomogeneities in the weld joint. The microinhomogeneity in the fusion zone, which could be characterized using electron backscatter diffraction studies, is a result of variations in morphology and formation of thermo-mechanically treated region across the weld pass interface. The considerable strength variation between the heat affected zone, base metal and fusion zone resulted in macro-inhomogeneity in the weld joint and introduced significant stress gradients which could be illustrated using finite element analysis. Both micro and macro-inhomogeneities have a synergistic role in determining the damage initiation location in multi-pass stainless steel weld joints subjected to creep.
1. Introduction Austenitic stainless steel is the prime structural material used for fabricating fast breeder reactor's structural components [1]. This material is a unanimous choice as it exhibits superior creep strength, adequate hot corrosion resistance and has good compatibility with liquid sodium [2]. As welding is the most widely resorted technique for fabricating huge components of the reactor, evaluating the creep strength of the weld joints is also as crucial as assessing the creep properties of the base material. Creep strength of a material is usually evaluated based on its ability to resist high temperature deformation as well as microstructural damage. Both deformation and microstructural damage influence the rupture life of the material. As components in nuclear reactors operate well within the allowable design stress limits, it is the time dependent damage in the material which is the predominant life limiting factor. Creep damage in austenitic stainless steel is usually attributed to nucleation and growth of cavities [3]. Creep cavitation reduces the effective load bearing capability of the component, which may in turn result in dimensional intolerance due to the consequent localised deformation. Elaborate assessment of the creep damage which results due to cavitation in the fusion zone, can offer new perspectives for any possible microstructural and design modifications. This in turn can help mitigate the influence of cavitation
⁎
Corresponding author. E-mail address: [email protected] (V.D. Vijayanand).
http://dx.doi.org/10.1016/j.msea.2017.07.057 Received 1 June 2017; Received in revised form 17 July 2017; Accepted 18 July 2017 Available online 19 July 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
induced damage on the failure of welded austenitic stainless steel components. In case of the austenitic stainless steel base metal, cavities during creep exposure usually initiate as a result of interactions between the grain boundary sliding phenomenon and chromium carbide precipitates [3]. For shorter durations of creep exposure when the carbides are finer in size, they pin grain boundaries effectively which obstructs sliding. Prolonged elevated temperature exposure results in significant coarsening of the carbides which render them as inefficient barriers for grain boundary sliding. This results in considerable grain boundary sliding resulting in nucleation of cavities at either triple points or at the chromium carbide precipitate interface [4]. In case of the fusion zone, regions containing delta ferrite exhibit creep cavitation [5]. A certain amount of delta ferrite is intentionally prescribed in austenitic steel fusion zone to prevent hot cracking [6]. The grain size in the fusion zone is coarser than the base metal [7] and has higher grain boundary facet length making the microstructure more susceptible to cavitation. As delta ferrite phase is a metastable phase with a relatively open body-centred cubic crystal structure when compared to the close packed face-centred cubic structure of austenite, it readily transforms into brittle intermetallic phases like sigma, chi, Laves. These transformations in addition to the precipitation of chromium carbide occur at shorter time durations when compared to the
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transformations occurring in the close-packed austenite phase. As the intermetallics and precipitates coarsen, their role in resisting grain boundary sliding diminishes gradually, resulting in nucleation of cavities at the interface [8]. Further, the presence of coarser grain size in the fusion zone results in higher facet length, which reduces the threshold stress required for cavity nucleation [9]. Thus it can be stated that in case of fusion zone the transformation of delta ferrite and in case of base metal the coarsening of carbides are the crucial phenomena which influence creep cavitation in each of these respective regions. It should be noted here that in a composite austenitic stainless steel weld joint which consists of base metal and fusion zone, the precipitation of intermetallic phases and chromium carbides from delta ferrite in the fusion zone occurs more rapidly at elevated temperature exposures when compared to the precipitation and subsequent coarsening of carbides in the austenitic base metal [10]. This results in an earlier onset of creep cavitation in the fusion zone when compared to the base metal region of the weld joint. Therefore, studies pertaining to damage evolution in the fusion zone of stainless steel are more essential for assessing the performance of the welded components. Unlike the base metal, microstructure of the fusion zone is highly heterogeneous, which makes studies on creep cavity nucleation more complicated. The primary source for the formation of heterogeneities is the laying of multiple passes and the associated thermal cycling which modifies the microstructure of the weld metal which is deposited by preceding passes. It has been reported that two distinct modifications occur in the weld metal due to the deposition of the subsequent pass. One of the modifications is associated with change of delta ferrite morphology and the other is with respect to the formation of an intrinsic thermo-mechanically strengthened region [11,12]. Studies on shielded metal arc welding (SMAW) joints have shown that deposition of subsequent pass modifies the structure of delta ferrite from vermicular to globular [11]. In order to comprehensively understand the prevalence of a thermo-mechanically region, investigations were undertaken on dual pass activated tungsten inert gas (A-TIG) welded joint. As A-TIG is an autogenous welding process, the influence of delta ferrite and its morphological changes was minimised. Using impression creep testing it was shown that the strength of the first laid pass was significantly higher than that of the pass which was laid subsequently. This study clearly established that the deposition of subsequent passes strengthen the previously deposited weld layers [13]. Cavitation in the fusion zone can be linked to the nature and distribution of the inhomogeneities. In this study, a thorough investigation is undertaken to evaluate the influence of various types of inhomogeneities on creep cavitation. Based on these results the possible reasons for creep cavitation to initiate at certain specific regions of the fusion zone have been critically examined.
Table 2 Welding parameters. Electrode size (mm) Current (A) Voltage (V) Travel speed (mm/min) Heat input (kJ/mm)
Fig. 1. Schematic of weld joint specimen.
schematic of the specimen is given in Fig. 1. Uniaxial creep tests were carried out at stress levels of 140 and 200 MPa at 923 K. Automated ball indentation (ABI) testing was carried out using a spherical indenter of 0.76 mm diameter. The indentor and was made of silicon carbide. ABI tests were performed at a temperature of 923 K and speed of the cross head was maintained at 0.008 mm s−1. Impression creep (IC) tests were carried out at 923 K using a flat bottomed cylindrical indenter of 1 mm diameter. For metallographic analysis, the weld joint specimens were polished by standard metallographic procedure up to a surface finish of 1 µm. These specimens were subsequently electrolytically etched using a solution containing 60% HNO3 and 40% H20 solution at 2 V for 30 s before observation. Micro-hardness measurements were taken using Vickers hardness tester employing a load of 200 gf. Specimens for electron back scatter diffraction (EBSD) were electro-polished before observation. EBSD data was acquired using Carl Zeiss Supra 55 FEGSEM. The FEG was subjected an acceleration voltage of 20 kV. The specimen was placed in a 70° pre-tilted holder with a working distance of 16 mm. The distribution of low and high angle grain boundaries was characterized using HKL-Channel 5 software. Finite element analysis was carried out using ABAQUS version 6.11 finite element solver. The four-node three dimensional linear tetrahedron element (C3D4) was used for meshing the geometry. The mesh size was refined until convergent values for the stress values were obtained.
3. Results and discussions 3.1. As-received microstructure The macrostructure of the weld joint is the as-welded condition is shown in Fig. 2. It can be seen from the figure that the fusion zone is composed of several layers. The fusion zone consisted of dendritic austenite grains with 3–5 vol% delta ferrite. Fig. 3(a and b) show the hardness variation across the X-Xʹ (base metal-HAZ-fusion zone-HAZbase metal) and Y-Yʹ (crown to root region of the fusion zone) lines indicated in Fig. 2. It can be seen that the hardness in the fusion zone and heat affected zone are substantially higher when compared to the base metal. The fusion zone is hardened due to the restraint imposed by the adjoining base metal and the solidification stresses which causes considerable increase in dislocation density in this dendritic microstructure [14]. It has been reported that the dislocation structure in the fusion zone of stainless steel can be compared to that of the heavily cold worked structures [15]. The HAZ in stainless steel weld joints is hardened due to the repeated thermal cycling generated during the multipass welding [16]. The increase in hardness from the crown to root region (Y-Yʹ) indicates that the root region is significantly hardened due to the deposition of subsequent passes. The dashed lines in Fig. 2
2. Experimental The chemical composition of the base material and electrode is given in Table 1. The details of the SMAW parameters are furnished in Table 2. The groove angle in the base plate was 10°. The weld joint specimen was composed of the three constituents the fusion zone in the central region, flanked by the HAZ as well as the base metal regions. A Table 1 Chemical composition range of 316LN SS base metal and 316(N) electrode in wt%.
316 LN SS base metal 316 (N) SS weld metal
C
Ni
Cr
Mo
Mn
Si
S
P
N
0.025
12.15
17.57
2.53
1.74
0.2
0.004
0.017
0.14
0.051
11.2
18.45
1.94
1.37
0.48
0.005
0.024
0.13
4 128 24 97 1.89 (each pass)
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The change in morphology was observed across all weld pass interfaces from the near-crown to the near-root regions. The transformation to globular morphology occurs due to the shape instability associated with the vermicular structure [17]. In addition to the morphological change which occurs in the globular region, deposition of subsequent passes results in the formation of a thermo-mechanically treated region. As a result, globular ferrite region exhibits better deformation resistance when compared to the vermicular region. The variations in morphology and formation of a thermo-mechanically treated region create microinhomogeneity within the fusion zone. Though the modifications in morphology were evident in both the near-crown and near-root regions, the extent of the strength mismatch across the weld pass interface could vary substantially. Strength gradients across the weld pass interface in the as-welded condition can be quantified using localised misorientation measurements employing EBSD studies. Fig. 5(a and b) shows the distribution of low angle grain boundaries (LAGB) and high angle grain boundaries (HAGB) at regions adjoining the weld pass interface taken from the near-crown and near-root regions. It can be seen from the Fig. 5c that the variation in LAGB distribution across the weld pass interface is more significant in the nearcrown region. The increase in LAGB is a signature of increased local strain gradients in the material [18]. The occurrence of less prominent gradient across the interface in the near-root pass region can be attributed to homogenisation of the dislocation substructure due to the laying of subsequent passes. Though the thermal cycle modifies the microstructure of the previously laid weld passes, it tends to accumulate more restraint in regions nearer to the root pass, when multiple weld layers are deposited subsequently. The increased restraint in the root pass results in higher hardness in this region when compared to the crown region (Fig. 3b). As the number of passes which are being laid on the previous pass reduces steadily from the root to the crown direction, this restraint reduces from the root to the crown region. As a consequence, the strength gradient across the weld pass interface is more pronounced near the crown region as it is subjected to fewer thermal cycles. Hence it can be concluded that the micro-inhomogeneity across the weld pass interface is more prominent in the crown region when compared to the root region.
Fig. 2. Macrostructure of the weld joint in the as-welded condition, the dashed line shows the region along which the creep specimen was extracted. The curved lines indicate the weld pass interfaces. The continuous lines X-X′ and Y-Y′ denote lines along which hardness measurements were taken.
Fig. 3. Hardness variation across the a) longitudinal (XX′) and b) transverse direction (YY′).
3.1.2. Macro-inhomogeneity From the hardness plot (Fig. 3a) taken across the various constituents of the weld joint, it is clear that there is a significant strength gradient in the weld joint specimen. This strength gradient results in macro inhomogeneity, generating complex stress distribution in the weld joint. The influence of fusion zone in generating stress gradients is similar to those generated due to presence of a mechanical notch. For this reason, the fusion zone is termed as a “metallurgical notch” in a weld joint specimen [19]. Miniature specimen techniques were used in this study to assess the mechanical properties of the fusion zone, HAZ and base metal independently to clearly establish the nature of stress
indicate the region from where the weld joint creep specimen was extracted. Two types of inhomogeneities were investigated within the fusion zone region bounded by the dashed lines. The region in the weld joint specimen adjacent to the crown is designated as near-crown region and the region near the root pass as near-root region in this study.
3.1.1. Micro-inhomogeneity During the deposition of the subsequent passes, the morphology of the delta ferrite changes from vermicular to globular (Fig. 4(a,b,c)).
Fig. 4. a) Microstructure across the weld pass interface, the dashed line represents the weld pass interface and the regions marked V and G indicate the vermicular and globular regions respectively b) magnified micrograph of vermicular region and c) magnified micrograph of globular region.
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Fig. 5. EBSD maps showing the distribution of grain boundaries in the a) near-crown and b) near-root region c) plot of the number fraction of boundaries in the respective regions. The low angle boundaries (2–10°, marked in red) and high angle boundaries (> 10°, marked in black). The green dashed line indicates the weld pass interface. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
distribution in the specimen. The automated ball indentation testing system was used to obtain tensile properties through a stress strain microprobe. This technique can be used to obtain the load (P)-depth of penetration (dp) plot of constituent regions in the weld joint using a spherical indentor. The equivalent true stress (σt) – plastic strain (εp) plot can be obtained using suitable conversion factors which are described below.
σt =
4P πd p2 δ
(1)
dp εp = 0.2 ⎛ ⎞ ⎝D⎠ ⎜
⎟
(2)
Here δ is the material dependent constraint factor which was taken as 3.1 in this study and D is the diameter of the indenter [20]. Fig. 6 shows the variation of yield strength and ultimate tensile strength in the fusion zone, HAZ and base metal regions of the weld joint obtained from the ABI technique tested at 923 K. It can be observed that the yield strength of the fusion zone is the highest when compared to the base metal and HAZ, however the ultimate tensile strength of this region is the lowest among the three regions. The fusion zone has considerably higher dislocation density as a result of solidification stress and constraint from the base metal. Further the dislocations can generate a tangled substructure during deformation whose motion can be more effectively blocked by interphase boundaries (delta ferrite-austenite boundary) when compared to the grain boundaries [21]. The UTS values are lower in the fusion zone due to the comparatively higher degree of localised plastic deformation of delta ferrite which has more slip systems at elevated temperatures [22]. The heat affected region has higher YS and UTS when compared to the base
Fig. 6. Variation of yield stress and ultimate tensile strength across the weld joint as measured by automated ball indentation technique.
metal due to the thermo-mechanical strengthening induced by repeated thermal cycling generated during the laying of multiple passes. Impression creep curve was used to independently evaluate the creep properties of the fusion zone, HAZ and base metal regions. The punching stress (σp) and steady state penetration rate of the indentor (vi) could be related to equivalent uniaxial stress (σu) and steady state . strain rate (ε u ) respectively, using the following equations.
σu = ασp
40
(3)
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Table 4 Creep test results. Base metal
Rupture life, hours Elongation, % Failure location
⎜
⎟
(4)
Where α and β are correlation factors dependent on the material. In the present investigation the values of α and β where taken as 0.33 and 1 respectively [23]. Fig. 7 shows the impression creep curves obtained using a punching stress of 681 MPa (uniaxial stress = 225 MPa). The curve clearly shows that both the overall depth of penetration and the steady state penetration rate was significantly lower for the fusion zone. The tangled dislocation substructure in this region results in comparatively higher resistance to creep deformation. Though the HAZ and base metal exhibited almost similar steady state penetration rates, the depths of penetration in these two regions were different. The HAZ region which was subjected to repeated thermal cycling, exhibited lower deformation rate and lower depths of penetration when compared to the unaffected base metal. The impression creep tests on each of the three constituent regions were repeated for various test levels to obtain the constants A and n of Norton's constitutive equation (ε˙ = Aσ n ) which establishes the dependence of the steady state creep rate (εε̇ ˙ ) with the applied stress (σ ). The values of A and n obtained from impression creep tests are given in Table 3. 3.2. Microstructure after creep testing A comparison between the creep test results of the weld joint and base metal is given in Table 4. Rupture life and elongation of the weld joint tested at 140 MPa were lower than that of the base metal. But at the stress level of 200 MPa, the rupture life of the weld joint was considerably higher than that of the base metal. The transformation of delta ferrite to brittle intermetallics and the associated creep cavitation in the fusion zone is the cause of pre-mature failure of the weld joint tested at 140 MPa. The failure of the weld joint tested at 200 MPa occurred at the fusion zone/HAZ interface due to strain incompatibility between these two microstructurally distinct regions. The rupture life of this weld joint was considerably short when compared to that of the weld joint tested at 140 MPa, thereby limiting the influence of creep cavitation in the fusion zone. The constraint provided by the fusion zone enhanced the rupture of the weld joint tested at 200 MPa when compared to the base metal. Though the failure location of the weld
Fusion zone −1
A(MPa h n
)
2.24 × 10 4.5
−15
HAZ 6.46 × 10 3.9
Base metal −13
200 MPa
140 MPa
175 MPa
9679
537
8060
1002
34
22.5
3.5 Within fusion zone
19 HAZ/fusion zone interface
3.2.1. Influence of micro-inhomogeneity Due to microstructural variations across the weld pass interface in the crown region, there is significant strain partitioning between the globular and vermicular delta ferrite regions under the influence of external stress. The strain partitioning due to the strength mismatch between the vermicular and globular ferrite region results in preferential nucleation of cavities in the vermicular region. The propagation of these cavities are however restricted at the weld pass interface. The globular delta ferrite region on the other side of the weld pass interface offers more resistance to propagation of cavities. The weld pass interface in the crown region obstructs the propagation of cavities. It can also be stated that in this region the interface serves as a boundary separating the cavity prone and cavity resistant regions. But in the root region the strength mismatch between the vermicular and globular delta ferrite is not significant. Therefore, under the influence of an external stress, strain partitioning is not prominent between the vermicular and globular ferrite regions in the root region. The absence of strain partitioning circumvents cavity propagation across the weld pass interfaces in the root region which renders it more resistant to creep cavitation.
Table 3 Values of Norton's constants A and n.
-n
140 MPa
joint varied with applied stress, considerable cavitation could be observed under both the test conditions. The macrostructure of the weld joint specimen tested at a stress level of 140 MPa at 923 K is shown in Fig. 8a. Fig. 8b shows a magnified region of the fusion zone near the weld pass interface. Preferential cavitation could be seen in the vermicular region of the weld pass interface. During creep exposure, the vermicular ferrite region near the weld pass interface is more susceptible to creep damage due to the favourable kinetics for transformation of delta ferrite to the brittle intermetallic phases [8]. The transformations are more sluggish in the globular morphology and this strengthened region offers more resistance to crack propagation and as a consequence, the cracks are blocked at the weld pass interface. The macrostructure also shows that in addition to the predominant crack which resulted in ultimate failure, there was another interlinked crack propagating to more than half of the specimen diameter. This suggests that the failure initiated in the near-crown region of the weld joint and propagated towards the root side. The macrostructure of the weld joint tested at higher stress level of 200 MPa is shown in Fig. 9. Since the failure of this joint occurred at the fusion zone/HAZ interface, damage distribution in the fusion zone was evident more explicitly. The density of cavities in this weld joint was quite high near the crown side and reduced steadily near the root side. Similar observation could be made for the joint tested at 140 MPa, where significant number of un-propagated cavities was observed in the crown region when compared to the root region. But there was a distinct variation in the location of the cavities with respect to distance from the weld centre line for the weld joint tested at 200 MPa. It could be clearly seen that the fusion zone cavitation in this weld joint occurred at regions adjacent to the HAZ interface, whereas for the weld joint tested at 140 MPa, the propagated cracks were almost along the weld centre line. The influence of micro and macro inhomogeneities on the cavity initiation and propagation is assessed in the following sections.
Fig. 7. Variation of depth of penetration with time for the three constituent regions as obtained from impression creep testing. . v ε u=⎛ i ⎞ βd ⎝ ⎠
Weld joint
1.05 × 10−12 3.9
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Fig. 8. a) Macrostructure of weld joint tested at 140 MPa/923 K, region A and B marked in the figure denote the crack initiation and final failure locations respectively, the arrows indicate the direction of the applied stress b) microstructure showing preferential cavitation in the same weld joint.
analysis. As Norton's law is only applicable in the secondary creep regime, the simulation was carried out for duration much shorter than the apparent secondary creep regime in the weld joint. This approach was adopted because from the creep curve of a composite weld joint it is not feasible to determine the relative contributions of the three constituent regions to the cumulative creep strain. Any possible localised microstructural or geometrical changes in any of the three regions would result in deviation from steady state behaviour. In order to offset any localised effect in the weld joint, the simulation was carried for duration of 6000 and 750 h, much lower than the rupture life of the specimen tested at 140 MPa and 200 MPa respectively. The variation of the stress tri-axiality factor obtained from the FE simulation was examined in this study due to its definitive role in influencing creep cavitation behaviour. Stress tri-axiality factor can be estimated using the following relationship.
Fig. 9. Macrostructure of weld joint tested at 200 MPa/923 K. The arrows indicate the direction of the applied stress.
3.2.2. Influence of macro-inhomogeneity The presence of microstructurally varying constituents results in significant stress gradients across the weld joint, which get amplified due to the presence of the groove angle. In order to study the evolution of the state of stress during creep, finite element analysis was performed on simulated weld joint geometry. One half of the composite weld joint specimen was modelled for the study and the simulation was performed under external stress levels of 140 and 200 MPa. The partitioning of the geometry into the fusion zone, heat affected zone and base metal was done based on the microstructure and hardness readings (Fig. 10). The four transverse lines indicated in the figure denote the paths along which the stress gradients were analysed. Path 1 was drawn along the weld centre line and paths 2, 3 and 4 were separated by a constant distance of 2 mm. Path 4 was closest to the fusion zone and HAZ interface. The tensile and creep properties used in the simulation were obtained by automated ball indentation and impression creep testing techniques. Norton's constitutive equation was used to simulate the plastic behaviour of the material under creep in the finite element
Stress Tri−axiality Factor =
σ1 + σ2 + σ3 3 × σvm
(5)
Where σ1, σ2, σ3 are the principal stress components and σvm is the vonMises stress. It has been well documented that both the von-Mises stress and the maximum principal stress influence nucleation and growth of cavities [24]. The tri-axiality factor balances the influence of both these factors. Fig. 11a and b show the distribution of stress tri-axiality factor for the weld joint geometries simulated at stress levels of 140 and 200 MPa. It could be seen from these figures that the presence of microstructurally varying constituent regions generate a significant variation in the state of stress across the weld joint. The influence of groove angle on the state of stress could be more clearly visualised in Fig. 12a and b. For the weld joint tested at both the stress levels, the tri-axial state of stress along transverse direction is almost constant near the weld centre line and starts to show a gradient with increase in distance from weld centre line. The variation in the tri-axiality factor in Paths 1, 2 and 3 was almost symmetric about the specimen's longitudinal axis. This trend could possibly be attributed to the relatively small groove angle (10°), which did not influence the variation of tri-axiality factor along these paths. Whereas for the path adjoining the interface (Path 4), the stress tri-axiality showed significant gradient with a minimum near the longitudinal axis of the specimen and increase in values along both the radial directions. It could also be seen that stress tri-axiality factor was marginally higher in the crown region when compared to the root region. Thus depending on the distance from the weld centre line, the presence of the groove angle can generate can generate stress gradients along the crown to root direction. In case of the weld joint tested at 140 MPa, the micro-inhomogeneity caused by the strength mismatch across the weld pass interface is more dominant in initiating creep cavitation in the crown region. The variation of stress tri-axiality factor resulting due to the
Fig. 10. Partitioned weld joint based on the microstructure and hardness. The red transverse lines in the figure indicate the paths along which the tri-aixiality factors are subsequently plotted. The blue longitudinal dashed line indicates the specimen's longitudinal axis. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
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Fig. 11. Contours of stress tri-axiality factor for weld joint simulated with applied stress of a) 140 MPa and b) 200 MPa.
comprising dimples which is a signature of ductile failure (Fig. 13a, b and c). The regions of the fractograph could be well correlated to the corresponding optical micrograph (Fig. 8a). Inter-dendritic failure can be linked to the propagation of cavities along the dendritic boundaries where prior transformed delta ferrite phase segregates during solidification. During creep exposure, the brittle intermetallics which transform from delta ferrite are also found segregated at the inter-dendritic boundaries. For the joint under the current investigation the cavity propagation occurred in the region marked A which corresponds to the near-crown region. With creep exposure, the propagation occurred into the specimen by possible inter-linking with other cavities nucleated in the bulk of the specimen. After significant propagation of the cracks the final failure of the specimen occurred in region B. Region B in the fractograph which showed evidences of ductile failure corresponds to the failure mode in the near-root pass region. The region adjacent to the root pass is strengthened significantly due to the laying of subsequent passes as a result it offered more resistance to inter-dendritic cavity propagation. The ductile failure in this region was a result of significant plastic deformation caused by stress ‘overload’. This is because the propagation of cavities causes significant reduction in load bearing cross sectional area and the final ductile failure is a result of significant plastic deformation, due to the associated stress concentration in the root pass region
groove angle could not have possibly influenced creep cavitation as the failure location was almost near the weld centre line. It can be seen that the distribution and values of tri-axiality factor in the fusion zone were similar for both the weld joint tested at two different stress levels. However, as the associated von-Mises stress (which is dependent on the applied stress) is much lower in case of the weld joint specimen tested at 140 MPa, the possibility of nucleation of cavities due to the geometrical constraint is inconsequential. Under the influence of a higher stress level of 200 MPa, the effects of macro-inhomogeneity are magnified causing preferred nucleation in the crown region of the fusion zone adjoining the HAZ interface. Even in the crown region, the microinhomogeneity due to the strength mismatch resulted in cavitation more specifically in the vermicular region. Therefore, in case of the weld joint tested at 200 MPa, both micro and macro-inhomogeneities played a synergistic role in initiating creep cavitation. It should also be noted in this weld joint that though the stress tri-axiality factor in the crown region was highest along path 4 (Fig. 12b), cavitation occurred only at a finite distance from the HAZ interface. This is because, only in these regions the grain boundaries oriented to almost 45°, which is the direction most favourable for stress assisted grain boundary sliding to occur [25]. It can also be inferred from this study that a sharper groove angle has the potential to cause significant variation in the stress tri-axiality factor across the transverse direction at shorter distances from the weld centre-line. This can cause significant increase in the tri-axiality factor in the crown region thereby catalysing the preferential nucleation of cavities in this region.
4. Conclusion Creep failure in stainless steel weld joints initiate in the crown region and propagate into the root region. The damage initiation in the crown region can be attributed to two inhomogeneities – a) micro-inhomogeneity caused due to the variation in strength gradient in the crown and the root region and b) macro-inhomogeneity due to the presence of a groove angle which changes fusion zone geometry from crown to root. The micro-inhomogeneity is the predominant factor causing preferential damage initiation in the weld joint at a stress level
3.3. Fractography The location of crack initiation site, crack propagation and final failure generated distinct features in the fractograph of the weld joint specimen tested at 140 MPa. The fractograph composed of a predominantly intergranular dendritic failure and a small region
Fig. 12. Variation of stress tri-aixiality factor taken along Paths 1, 2, 3, 4 for weld joint simulated with applied stress of a) 140 MPa and b) 200 MPa.
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Fig. 13. Fractograph of specimen failed at 140 MPa, region A and B marked in the figure denote the crack initiation and final failure locations respectively, b) magnified image of region A and c) magnified image of region B. [9] A.G. Evans, Acta Metall. 28 (1980) 1155–1163. [10] J.J. Smith, R.A. Farrar, Int. Mater. Rev. 38 (1993) 25–51. [11] V.D. Vijayanand, K. Laha, P. Parameswaran, V. Ganesan, M.D. Mathew, Mater. Sci. Eng. A 607 (2014) 138–144. [12] V.D. Vijayanand, J. Ganesh Kumar, P.K. Parida, V. Ganesan, K. Laha, Metall. Mater. Trans. A 48 (2017) 706–721. [13] V.D. Vijayanand, M. Vasudevan, V. Ganesan, P. Parameswaran, K. Laha, A.K. Bhaduri, Metall. Mater. Trans. A 47 (2016) 2804–2814. [14] T.G. Gooch, J. Honeycombe, Weld. J. 59 (1980) 233–241. [15] M. Valsan, D. Sundararaman, K.B.S. Rao, S.L. Mannan, Metall. Mater. Trans. A 26 (1995) 1207–1220. [16] C.F. Etienne, D.V. Rossum, F. Roode, in: Proceedings of the International Conference on Engineering Aspects of Creep, University of Sheffield, 2, 1980, pp. 113–121. [17] S.A. David, Weld. J. 60 (1981) 63–71. [18] S.I. Wright, M.M. Nowell, D.P. Field, Microsc. Microanal. 17 (2011) 316–329. [19] M.J. Manjoine, Weld. J. 61 (1982) 50–57. [20] J. Ganesh Kumar, V.D. Vijayanand, M. Nandagopal, K. Laha, Mater. High Temp. 32 (2015) 619–626. [21] E. Werner, H.P. Stuwe, Mater. Sci. Eng. A 68 (1984 & ;;85) 175–182. [22] C.R. Weinberger, B.L. Boyce, C.C. Battaile, Int. Mater. Rev. 58 (2013) 296–314. [23] Naveena, V.D. Vijayanand, V. Ganesan, K. Laha, M.D. Mathew, Mater. Sci. Eng. A 552 (2012) 112–118. [24] T.L. Sham, A. Needleman, Acta Metall. 31 (1993) 919–926. [25] R.C. Gifkins, A. Gittins, R.L. Bell, T.G. Langdon, J. Mater. Sci. 3 (1968) 306–313.
of 140 MPa. In addition to this reason, the variation in macro-inhomogeneity could also influence creep cavitation for the weld joint tested at a higher stress level of 200 MPa. Though cavitation in the fusion zone of stainless steel weld joints occurs in the crown region, depending on the applied stress and groove angle its location from the weld centre line can vary substantially. References [1] S.L. Mannan, S.C. Chetal, Baldev Raj, S.B. Bhoje, Trans. Indian Inst. Met. 235 (2005) 155–178. [2] Baldev Raj, P. Chellapandi, P.R. Vasudeva Rao, Sodium Fast Reactors with Closed Fuel Cycles, CRC Press, Taylor and Francis Group, Florida, 2015, pp. 86–90. [3] I.W. Chen, A.S. Argon, Acta Metall. 29 (1981) 1321–1333. [4] G. Sasikala, M.D. Mathew, K.B.S. Rao, S.L. Mannan, Metall. Mater. Trans. A 31 (2000) 1175–1185. [5] C.A.P. Horton, P. Marshall, R.G. Thomas, Mechanical Behavior and Nuclear Applications of Stainless Steel at Elevated Temperatures, The Metals Society, London, 1982, pp. 66–72. [6] K.H. Lo, C.H. Shek, J.K.I. Lai, Mater. Sci. Eng. R. 65 (2009) 39–104. [7] J.C. Lippold, W.F. Savage, Weld. J. 59 (1980) 48–58. [8] B.A. Senior, J. Mater. Sci. 25 (1990) 45–53.
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