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Microstructural evolution during creep of 316LN stainless steel multi-pass weld joints
Materials Science & Engineering A 607 (2014) 138–144
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructural evolution during creep of 316LN stainless steel multi-pass weld joints V.D. Vijayanand n, K. Laha, P. Parameswaran, V. Ganesan, M.D. Mathew Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India
art ic l e i nf o
a b s t r a c t
Article history: Received 29 January 2014 Received in revised form 28 March 2014 Accepted 31 March 2014 Available online 8 April 2014
Creep rupture behaviour of 316LN austenitic stainless steel weld joint fabricated by multi-pass shielded metal arc welding process has been studied at 923 K over a stress range of 120–225 MPa. The weld joint exhibited inferior creep rupture strength than the base metal and extensive creep cavitation in weld metal led to the premature failure. Weld metal microstructure was found highly inhomogeneous having different morphologies of delta (δ)-ferrite. The δ-ferrite transformed into intermetallic phases on creep exposure and the creep cavitation was associated with the intermetallic phases. Extensive creep cavitation in weld metal was confined to regions containing δ-ferrite with vermicular morphology. The region near the weld pass interface having globular δ-ferrite was less susceptible to creep cavitation. The globular δ-ferrite region possessed higher hardness than the vermicular δ-ferrite region. High dislocation density was observed in the globular δ-ferrite region which is a consequence of the microstructural modification by the heat input from the successive weld passes. The strength inhomogeneity between the globular and vermicular δ-ferrite regions together with the transformation of δ-ferrite into intermetallic phases produced pronounced creep cavitation in the vermicular δ-ferrite region which led to premature failure of the stainless steel weld joint. & 2014 Elsevier B.V. All rights reserved.
Keywords: Electron micrography Austenitic stainless steel Delta-ferrite morphology Multi-pass weld joint Intermetallic precipitation Creep cavitation
1. Introduction Liquid metal cooled fast breeder reactors employ 316LN austenitic stainless steels (SS) for its high temperature structural components [1]. It has been shown that addition of nitrogen improves the tensile, creep and low cycle fatigue strength of the steel [2–4]. Fast reactor components are fabricated essentially by fusion welding and evaluation of mechanical behaviour of 316LN SS weld joints is of paramount importance. At high temperatures, the mechanical strength of the base metal and the weld joints are not similar. Since weld joints are regions of geometric discontinuity and exhibit inhomogeneous microstructure, they exhibit inferior creep rupture strength when compared to the corresponding base metal. The δ ferrite which is incorporated in austenitic weld metal by intentional compositional adjustments for improving the hot cracking resistance during welding is the main source for microstructural instability during creep exposure, which influences the creep rupture properties of the weld joint appreciably. The content of δ-ferrite specified for hot corrosion resistance is 3–7 vol% [5]. The upper limit is fixed in view of preventing the precipitation of excess brittle intermetallics during elevated temperature exposure [6,7]. Careful methodology is adopted
n
Corresponding author. Tel.: þ 91 44 27480500x22216; fax: þ91 44 27480075. E-mail address: [email protected] (V.D. Vijayanand).
http://dx.doi.org/10.1016/j.msea.2014.03.138 0921-5093/& 2014 Elsevier B.V. All rights reserved.
to obtain a microstructure distribution containing a uniform distribution of δ-ferrite in the specified range; nevertheless multi-pass welding of especially thick components generates a heterogeneous microstructure containing δ-ferrite of different morphologies in the weld metal. Although several studies have been carried out on creep cavitation and rupture behaviour of austenitic stainless steel weld joints, very little information is available in the open literature which discusses the effects of heterogeneity in the microstructure of multi-pass weld joints [8] and its implication on creep cavitation behaviour [9,10]. The current investigation outlines creep rupture behaviour of fusion welded 316LN SS weld joints. Special emphasis has been laid on the effect of variation in the microstructure of weld metal brought about by multi-pass welding and its effect on elevated temperature creep cavitation and rupture of the weld joint.
2. Experimental The chemical composition of the steel and the electrode is shown in Table 1. The 316LN SS weld pads of 500 250 20 mm3 size were fabricated by employing multi-pass shielded metal arc welding process employing matching welding consumable. The weld pads
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Table 1 Chemical composition of 316LN SS base metal and 316(N) electrode in wt%. Material
C
Ni
Cr
Mo
Mn
Si
S
P
N
316 LN SS base metal 316 (N) SS weld metal
0.025 0.052
12.15 11.5
17.57 18.6
2.53 2.2
1.74 1.74
0.2 0.64
0.004 0.007
0.017 0.022
0.14 0.1
Fig. 2. Schematic showing the representative regions where EDS analysis and TEM samples were extracted.
Fig. 1. Schematic of weld pad showing locations where base metal and weld joint samples taken.
Table 2 Weld pad details. Process
Shielded metal arc
Joint design Inter-pass temperature Flux Current Voltage Travel speed Number of passes
Single V groove 423 K Basic coated 150 A 25 V 3 mm/s 32
were qualified by radiographic examination for their soundness. Specimens with 50 mm gauge length and 10 mm gauge diameter were used for the creep tests. Cross weld joint specimens consisting of weld metal and base metal were fabricated from the weld pads with gauge length parallel to the rolling direction of the plate. Fig. 1 shows the schematic of the weld pad and Table 2 gives the details of the welding parameters. Constant load creep rupture tests were carried out at 923 K at stresses of 120, 140, 175, 200 and 225 MPa. Ferritescope was used to measure δ-ferrite content in the weld metal both before and after creep tests. Hardness measurements were carried out using a Vickers hardness tester employing a load of 100 gf. Microstructural studies were carried out using optical and scanning electron microscopes (SEM). Immersion etching in boiling Murakkami etchant (10 g potassium ferric cyanide, 10 g potassium hydroxide and 100 ml water) was used to reveal the δ-ferrite and the transformed phases. The estimation of elemental distribution adjacent to weld pass interface was carried out using energy dispersive x-ray spectrum (EDS) analyser attached to the SEM. Two representative regions of weld metal as shown in Fig. 2 were selected for elemental investigation. The regions were at a distance of 3 mm from the failure location and around 150 μm away from the weld pass interface. Data was collected from 5 fields of view and at least 20 precipitates were analysed in each field of view for each representative location. Thin foil examination of different morphologies of weld
Fig. 3. Macrostructure of the as weld material showing X–X0 and Y–Y0 sections where hardness values were taken, the dashed line shows the region from which the creep specimen was extracted.
metal was carried out using transmission electron microscopy (TEM). The samples were extracted from similar locations where EDS analysis was carried out.
3. Results and discussion 3.1. Microstructure of the weld joint in as-welded condition Macrostructure of the weld joint in as-received condition is shown in Fig. 3, indicating the regions from where the creep specimen was extracted along with the sections where hardness values were obtained. The weld joint broadly consisted of as-deposited weld metal, heat affected zones on both sides of deposited weld metal and the unaffected base metal. Microstructure across the weld metal was highly inhomogeneous with different morphologies of δ-ferrite. Two distinct morphologies of δ-ferrite namely vermicular and globular were observed. In vermicular ferrite region, the δ-ferrite was interconnected and presented in the interdendritic regions of austenite (Fig. 4(a)). This morphology was found in regions which had not been affected by weld thermal cycles due to subsequent weld pass. It constitutes major portion of the deposited weld metal. Ferrite of globular morphology was present in region near to the weld pass interface over which the subsequent weld pass was laid, in other
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Fig. 4. Microstructure of the as welded joint showing δ-ferrite with (a) vermicular and (b) globular morphologies.
Fig. 6. Variation of hardness along a weld pass interface. Fig. 5. Variation of hardness along X–X0 AND Y–Y0 sections.
words it can be said that this region is the ‘heat affected zone’ in the weld metal (Fig. 4(b)). The spread of this region was up to 700 μm from the weld pass interface. The variations of hardness across crown to root (Y–Y0 section in Fig. 3) and across base metal-heat affected zone to weld metal-heat affected zone-base metal (X–X0 section in Fig. 3) of the weld joint are shown in Fig. 5. Though hardness increased from crown to root of the weld metal (in the Y–Y0 section), there was appreciable scatter in the data. The weld metal in general exhibited higher hardness than the base metal due to its as-cast structure. The increase in hardness from crown to root of the joint can be attributed to the increased dislocation density generated by weld thermal cycles of the multipass welding. The appreciable scattering in hardness is considered due to the microstructural inhomogeneity across each weld pass. The hardness variation between successive weld passes is shown in Fig. 6. The globular δ-ferrite region which had formed as a result of thermal cycle produced by the subsequent weld pass, exhibited higher hardness than that of the vermicular ferrite region. There was similar variation in hardness of the weld metal across X–X0 section as in Y–Y0 . It is however important to note that the heat affected zone in the base metal exhibited higher hardness than the base metal, this again can be attributed to increased dislocation density produced by weld thermal cycle [11–13]. TEM micrograph of the vermicular and globular ferrite regions of the weld metal are shown in Fig. 7(a) and (b) respectively. The interconnected δ-ferrite, as observed from the optical micrograph (Fig. 3(a)), could be clearly seen along with free dislocations in both δ-ferrite and austenite in the vermicular ferrite region.
The substructure in the globular regions consisted of dislocation forest structure and isolated δ-ferrite with a globular morphology (Fig. 7(b)), giving rise to higher hardness (Fig. 6). As in case of the heat affected region of 316LN SS weld joints, the dislocation density was higher in this region subjected weld thermal cycle by the subsequent weld thermal cycle. Although through TEM analysis the spread of this region consisting of dislocation forest structure could not be established, hardness measurements indicated that this region matched with the region of globular ferrite morphology. 3.2. Creep rupture properties A comparison between creep rupture lives of the base metal and cross weld joint, creep tested at 923 K over at stresses in the range 120–225 MPa, is shown in Fig. 8. The rupture life of weld joints is substantially lower than that of the base metal, especially at relatively lower applied stresses. The difference in creep rupture strength between the base metal and weld joint enhanced with increase in rupture life. Failure of all the weld joints occurred in the weld metal region. The microstructural instability of δ-ferrite contained in weld region, as discussed subsequently, is considered for the reduction in creep rupture life of weld joint than the base metal. The variations of creep rupture ductility with creep rupture life of both the base metal and the weld joint are shown in Fig. 9. Creep ductility decreased with creep exposure for the weld joints whereas it increased for the base metal. The gradual reduction in creep rupture ductility in the weld joint can be attributed to the nucleation and growth of creep cavities associated with the
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Fig. 7. TEM micrograph of the as-weld material showing (a) vermicular interconnected δ-ferrite and (b) globular δ-ferrite with higher density of dislocations.
Fig. 10. Variation of residual δ-ferrite with rupture life. Fig. 8. Variation of rupture life with applied stress.
intermetallics. The base metal is more resistant to cavitation after longer creep exposure because the carbides in the grain boundaries coarsen and they offer very less resistance to grain boundary sliding. The residual δ-ferrite content steadily reduced with creep exposures and it could be ascertained that for the weld joint with longest rupture the transformation to brittle intermetallics is almost complete (Fig. 10). 3.3. Microstructural modification of weld metal on creep exposure
Fig. 9. Variation of rupture elongation with rupture life.
The elevated temperature transformation of δ-ferrite into carbides and intermetallics has been widely studied in austenitic stainless steel welds [14,15]. The transformation commences with the precipitation of M23C6 at the austenite/δ-ferrite interface. Prolonged exposure can lead to the precipitation of intermetallic phases such as s, χ and η. It has been argued that the amounts of C and N in the weld metal have definitive roles in the kinetics of precipitation of these intermetallics. While N inhibits the formation of all the aforesaid intermetallic phases, C on the other hand restricts the precipitation of s and η phases, but not the χ phase. This is because C is soluble in χ and the precipitation of this phase can happen simultaneously along with M23C6 [16]. Though precipitation of intermetallics can occur in austenite also, its kinetics is rather slow. The δ-ferrite has higher
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solubility for Cr, Mo which are the main constituents of the intermetallics. Further, the diffusion of these elements is higher in δ-ferrite when compared to austenite, because of its more open BCC structure. During creep exposure, the precipitation of the phases is enhanced than in thermal aging due to the stress induced diffusion of the elements [17]. In light of the above discussion it is clear that the essential factors for the precipitation of carbides are the diffusion of Cr and Mo. The formation of M23C6 depends on the diffusion of Cr, whereas diffusion of both Cr and Mo is essential for the precipitation of intermetallics [18]. It was also pointed out that continuous network of δ-ferrite results in higher δ-ferrite/austenite interface boundaries which act as nucleation sites for carbides and intermetallics. In the present investigation, the globular morphology consisted of isolated δ-ferrite which lowers interface boundaries. Thus the transformation of δ-ferrite into carbides and intermetallics is delayed in regions where delta ferrite exhibited globular morphology. Fig. 11 shows the variations of contents of Cr and Mo in the transformed phases with creep exposure, estimated by EDS attached with SEM, in both the globular and vermicular region of the weld metal. It can be clearly seen that the enrichment of Cr and Mo in the phases was higher in the vermicular regions when compared to the globular regions. Hence, it can be ascertained that the transformation of δ-ferrite into carbides and intermetallics could readily occur in the vermicular region than that in globular region. It should be also noted that multi-pass welding results in variations of the initial delta ferrite content in the globular and vermicular regions [9], and this can also influence the transformation kinetics.
3.4. Dislocation substructure evolution during creep The evolution of dislocation substructure was studied in detail for one short term test (applied stress 175 MPa, rupture life 1300 h) and one long term test (applied stress 120 MPa, rupture life 7820 h). The weld joint tested at 175 MPa showed dislocation pile ups and slip bands in the vermicular region (Fig. 12(a) and (b)). However in the globular region, subgrains could be seen (Fig. 12(c)). Precipitation of intermetallics was limited in this test condition. Fig. 13(a) shows the dislocation substructure in the vermicular region of the weld joint tested at 120 MPa. The cell structure which is a characteristic of recovery during creep of alloys with low stacking faults is clearly visible along with coarsened intermetallics. EDS spectra on these intermetallics revealed that they were Fe–Cr–Mo and Mo–Fe rich suggesting that they could be s and η respectively (Fig. 14). The TEM micrograph taken in the globular region also showed intermetallics and subgrains (Fig. 13(b)). However, the dislocation distribution inside the subgrains was higher in this region when compared to those in the vermicular region. The evolution of dislocation substructure during creep of austenitic steel welds is well documented [19,20]. At high stress levels, in a well annealed microstructure such as the one in the vermicular region of the weld metal in the current investigation – planar slip is the principal deformation mechanism and the matrix resistance to plastic deformation is quite negligible. A cold worked structure with higher dislocation density such as the one in the globular region, subgrains form which is the signature of recovery during creep deformation. Although prolonged creep exposures (creep exposure at lower stress levels) produced subgrains in the in both vermicular and globular regions, there were some remnant dislocations within the subgrains of globular regions. This suggests that the formation of subgrains is not complete in this region. The excessive size of the intermetallics would not have caused any significant interactions with the dislocations in both the regions [19]. Hence it can be assumed that the initial dislocation substructure is the decisive factor driving recovery in both the regions of the weld at this stress level. The initial dislocations density was less and there were free dislocations in the vermicular regions when compared to a more tangled dislocation structure in the globular region. The applied stress in this case was not high enough for recovery mechanisms such as dislocation climb to operate; hence formation of sub grains was delayed in the dislocation rich globular region. 3.5. Creep damage and failure.
Fig. 11. Variation of Cr and Mo in transformed regions with rupture life.
Microstructure of the weld joint tested at 175 MPa showed intragranular cracking in the vermicular regions near the weld
Fig. 12. (a) Dislocation pile up, (b) slip bands in vermicular region and (c) subgrains formation in globular region after creep testing at 175 MPa.
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Fig. 13. TEM micrographs of weld joint tested at 120 MPa (rupture life 7820 h) (a) vermicular region consisting subgrains and precipitate A with Fe–Cr–Mo and B with Fe–Mo enrichment and (b) globular region with subgrains.
Fig. 14. EDS spectra of (a) precipitate A showing enrichment of Fe–Cr–Mo and (b) precipitate B showing Fe–Mo enrichment.
Fig. 15. (a) Intergranular cracking occurred for at 175 MPa and (b) creep cavitation occurred along the intermetallic/matrix after creep testing at 120 MPa.
Fig. 16. (a) Fractograph of weld joint tested at 175 MPa, (b) region I showing interdendritic failure and (c) region D showing ductile failure.
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Fig. 17. (a) Fractograph of weld joint tested at 120 MPa, (b) region B showing cracks in brittle transformed intermetallics and (c) region D showing ductile failure.
pass interface (Fig. 15(a)) whereas regions with globular morphology did not exhibit any cracking. As discussed earlier there is a significant difference in the dislocation substructure in the two regions which is also reflected in the hardness values reported earlier. This results in stress concentration in the regions very near the interface, as a consequence there is tendency for intergranular cracking to occur in comparatively softer region, which in the present context refers to the vermicular region. Fig. 15(b) shows the SEM micrograph in the vermicular region of the weld joint tested at 120 MPa. Interconnected cavitation along the transformed regions is clearly visible in this region. As pointed out earlier, the morphology of δ-ferrite in these regions is favourable for transformation kinetics of in this region, and is more prone to creep cavitation. The origin of creep damage in the intermetallic/ matrix interface has been discussed by many researchers in austenitic stainless steel weld metal [21–23]. Figs. 16 and 17 show the fractographs weld joints tested at 175 and 120 MPa respectively. The aforementioned cracks which formed in the intergranular regions for the weld joint tested at 175 MPa, manifested as fibrous region (Fig. 16(b)). Whereas the weld joint tested at 120 MPa showed regions of extensive cracking in the transformed brittle intermetallics (Fig. 17(b)). Both the fractographs contained ductile features which could have possibly formed due to failure during ‘overload’ in the globular regions of the weld metal (Figs. 16 (c) and 17(c)). But it could be clearly seen that the extent of ductile region was more in the weld joint crept for 1300 h (applied stress 175 MPa) due to the incomplete transformation of δ-ferrite. 4. Conclusion Creep rupture properties of multi-pass SMAW 316LN SS was studied in the stress levels of 120–225 MPa at 923 K and the following conclusions have been drawn. 1. Depending on the creep testing conditions, the failure of the weld joints was attributed to two different phenomenons occurring due to the variations in morphology and dislocation substructure evolved due to multi-pass welding. 2. At higher stress levels the difference in dislocation substructure in the globular and vermicular regions led to stress concentration resulting in cracking of the softer vermicular region. 3. At low stress levels, creep cavitation was confined in the vermicular region due to the presence of transformed intermetallics and the strength inhomogeneity with the adjoining globular region. 4. Multi-pass welding produced regions of high dislocation density in the region consisting globular ferrite, rendering it more
resistant to creep deformation at high stress levels. At lower stress levels, where the failure was attributed to intermetallic creep cavitation, the transformation of δ-ferrite was delayed due to its globular morphology.
Acknowledgement The authors thank Dr. P.R. Vasudeva Rao, Director, Indira Gandhi Centre for Atomic Research, Kalpakkam and Dr. T. Jayakumar, Director, Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam, India for their continual support and guidance. References [1] S.L. Mannan, S.C. Chetal, Baldev Raj, S.B. Bhoje, Trans. Indian Inst. Met. 56 (2003) 155–178. [2] M.D. Mathew, K. Laha, V. Ganesan, Mater. Sci. Eng. A 535 (2012) 76–83. [3] V. Ganesan, M.D. Mathew, K. Bhanu Sankara Rao, Mater. Sci. Technol. 25 (2009) 614–618. [4] G.V. Prasad Reddy, R. Sandhya, K.B.S. Rao, S. Sankaran, Proc. Eng. 2 (2010) 2181–2188. [5] F.C. Hull, Weld. J. 60 (1967) 399–409. [6] J.J. Smith, R.A. Farrar, Int. Mater. Rev. 38 (1993) 25–51. [7] S.R. Keown, R.G. Thomas, Met. Sci. 15 (1981) 386–/392. [8] Y. Song, T.N Baker, N.A McPherson, Mater. Sci. Eng. A 212 (1996) 228–234. [9] C.A.P. Horton, J.K. Lai, Met. Sci. 14 (1980) 502–505. [10] Q. Auzoux, L. Allais, C. Caes, I. Monnet, A.F. Gourgues, A. Pineau, J. Nucl. Mater. 400 (2010) 127–137. [11] Naveena, V.D. Vijayanand, V. Ganesan, K. Laha, M.D. Mathew, Mater. Sci. Technol. (2013), http://dx.doi.org/10.1179/1743284713Y.0000000438. [12] P. Marshall, Austenitic Stainless Steels: Microstructure and Mechanical Properties, Elsevier Applied Science Publishers, London and New York, 1984. [13] C.F Etienne, D. Van Rossum, F. Roode, Proceedings of the International Conference on Engineering Aspects of Creep, September 1980, University of Sheffield, U.K., vol. 2, pp. 113–121. [14] K.H. Lo, C.H. Shek, J.K.I. Lai, Mater. Sci. Eng. R 65 (2009) 39–104. [15] G. Sasikala, M.D. Mathew, K.B.S. Rao, S.L. Mannan, Met. Mater. Trans. A 31 (2000) 1175–1185. [16] E. Folkhard, Welding Metallurgy of Stainless Steels, Springer Verlag, Wein, New York, 1986. [17] M.D. Mathew, G. Sasikala, T.P.S. Gill, S.L. Mannan, P. Rodriguez, Mater. Sci. Technol. 10 (1994) 1104–1108. [18] G. Sasikala, S.K. Ray, S.L. Mannan, Mater. Sci. Eng. A 359 (2003) 86–90. [19] M.D. Mathew, M. Sundararaman, S.L. Mannan, Mater. Trans. 38 (1997) 37–42. [20] J.R. Foulds, J. Monteff, Weld. J. 61 (1982) 189–196. [21] B.A. Senior, J. Mater. Sci. 25 (1990) 45–53. [22] G. Sasikala, S.K. Ray, S.L. Mannan, Acta Mater. 52 (2004) 5677–5686. [23] T. Shakhivel, M. Vasudevan, K. Laha, P. Parameswaran, K.S. Chandravathi, M.D. Mathew, A.K. Bhaduri, Mater. Sci. Eng. A 528 (2011) 6971–6980.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructural evolution during creep of 316LN stainless steel multi-pass weld joints V.D. Vijayanand n, K. Laha, P. Parameswaran, V. Ganesan, M.D. Mathew Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India
art ic l e i nf o
a b s t r a c t
Article history: Received 29 January 2014 Received in revised form 28 March 2014 Accepted 31 March 2014 Available online 8 April 2014
Creep rupture behaviour of 316LN austenitic stainless steel weld joint fabricated by multi-pass shielded metal arc welding process has been studied at 923 K over a stress range of 120–225 MPa. The weld joint exhibited inferior creep rupture strength than the base metal and extensive creep cavitation in weld metal led to the premature failure. Weld metal microstructure was found highly inhomogeneous having different morphologies of delta (δ)-ferrite. The δ-ferrite transformed into intermetallic phases on creep exposure and the creep cavitation was associated with the intermetallic phases. Extensive creep cavitation in weld metal was confined to regions containing δ-ferrite with vermicular morphology. The region near the weld pass interface having globular δ-ferrite was less susceptible to creep cavitation. The globular δ-ferrite region possessed higher hardness than the vermicular δ-ferrite region. High dislocation density was observed in the globular δ-ferrite region which is a consequence of the microstructural modification by the heat input from the successive weld passes. The strength inhomogeneity between the globular and vermicular δ-ferrite regions together with the transformation of δ-ferrite into intermetallic phases produced pronounced creep cavitation in the vermicular δ-ferrite region which led to premature failure of the stainless steel weld joint. & 2014 Elsevier B.V. All rights reserved.
Keywords: Electron micrography Austenitic stainless steel Delta-ferrite morphology Multi-pass weld joint Intermetallic precipitation Creep cavitation
1. Introduction Liquid metal cooled fast breeder reactors employ 316LN austenitic stainless steels (SS) for its high temperature structural components [1]. It has been shown that addition of nitrogen improves the tensile, creep and low cycle fatigue strength of the steel [2–4]. Fast reactor components are fabricated essentially by fusion welding and evaluation of mechanical behaviour of 316LN SS weld joints is of paramount importance. At high temperatures, the mechanical strength of the base metal and the weld joints are not similar. Since weld joints are regions of geometric discontinuity and exhibit inhomogeneous microstructure, they exhibit inferior creep rupture strength when compared to the corresponding base metal. The δ ferrite which is incorporated in austenitic weld metal by intentional compositional adjustments for improving the hot cracking resistance during welding is the main source for microstructural instability during creep exposure, which influences the creep rupture properties of the weld joint appreciably. The content of δ-ferrite specified for hot corrosion resistance is 3–7 vol% [5]. The upper limit is fixed in view of preventing the precipitation of excess brittle intermetallics during elevated temperature exposure [6,7]. Careful methodology is adopted
n
Corresponding author. Tel.: þ 91 44 27480500x22216; fax: þ91 44 27480075. E-mail address: [email protected] (V.D. Vijayanand).
http://dx.doi.org/10.1016/j.msea.2014.03.138 0921-5093/& 2014 Elsevier B.V. All rights reserved.
to obtain a microstructure distribution containing a uniform distribution of δ-ferrite in the specified range; nevertheless multi-pass welding of especially thick components generates a heterogeneous microstructure containing δ-ferrite of different morphologies in the weld metal. Although several studies have been carried out on creep cavitation and rupture behaviour of austenitic stainless steel weld joints, very little information is available in the open literature which discusses the effects of heterogeneity in the microstructure of multi-pass weld joints [8] and its implication on creep cavitation behaviour [9,10]. The current investigation outlines creep rupture behaviour of fusion welded 316LN SS weld joints. Special emphasis has been laid on the effect of variation in the microstructure of weld metal brought about by multi-pass welding and its effect on elevated temperature creep cavitation and rupture of the weld joint.
2. Experimental The chemical composition of the steel and the electrode is shown in Table 1. The 316LN SS weld pads of 500 250 20 mm3 size were fabricated by employing multi-pass shielded metal arc welding process employing matching welding consumable. The weld pads
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Table 1 Chemical composition of 316LN SS base metal and 316(N) electrode in wt%. Material
C
Ni
Cr
Mo
Mn
Si
S
P
N
316 LN SS base metal 316 (N) SS weld metal
0.025 0.052
12.15 11.5
17.57 18.6
2.53 2.2
1.74 1.74
0.2 0.64
0.004 0.007
0.017 0.022
0.14 0.1
Fig. 2. Schematic showing the representative regions where EDS analysis and TEM samples were extracted.
Fig. 1. Schematic of weld pad showing locations where base metal and weld joint samples taken.
Table 2 Weld pad details. Process
Shielded metal arc
Joint design Inter-pass temperature Flux Current Voltage Travel speed Number of passes
Single V groove 423 K Basic coated 150 A 25 V 3 mm/s 32
were qualified by radiographic examination for their soundness. Specimens with 50 mm gauge length and 10 mm gauge diameter were used for the creep tests. Cross weld joint specimens consisting of weld metal and base metal were fabricated from the weld pads with gauge length parallel to the rolling direction of the plate. Fig. 1 shows the schematic of the weld pad and Table 2 gives the details of the welding parameters. Constant load creep rupture tests were carried out at 923 K at stresses of 120, 140, 175, 200 and 225 MPa. Ferritescope was used to measure δ-ferrite content in the weld metal both before and after creep tests. Hardness measurements were carried out using a Vickers hardness tester employing a load of 100 gf. Microstructural studies were carried out using optical and scanning electron microscopes (SEM). Immersion etching in boiling Murakkami etchant (10 g potassium ferric cyanide, 10 g potassium hydroxide and 100 ml water) was used to reveal the δ-ferrite and the transformed phases. The estimation of elemental distribution adjacent to weld pass interface was carried out using energy dispersive x-ray spectrum (EDS) analyser attached to the SEM. Two representative regions of weld metal as shown in Fig. 2 were selected for elemental investigation. The regions were at a distance of 3 mm from the failure location and around 150 μm away from the weld pass interface. Data was collected from 5 fields of view and at least 20 precipitates were analysed in each field of view for each representative location. Thin foil examination of different morphologies of weld
Fig. 3. Macrostructure of the as weld material showing X–X0 and Y–Y0 sections where hardness values were taken, the dashed line shows the region from which the creep specimen was extracted.
metal was carried out using transmission electron microscopy (TEM). The samples were extracted from similar locations where EDS analysis was carried out.
3. Results and discussion 3.1. Microstructure of the weld joint in as-welded condition Macrostructure of the weld joint in as-received condition is shown in Fig. 3, indicating the regions from where the creep specimen was extracted along with the sections where hardness values were obtained. The weld joint broadly consisted of as-deposited weld metal, heat affected zones on both sides of deposited weld metal and the unaffected base metal. Microstructure across the weld metal was highly inhomogeneous with different morphologies of δ-ferrite. Two distinct morphologies of δ-ferrite namely vermicular and globular were observed. In vermicular ferrite region, the δ-ferrite was interconnected and presented in the interdendritic regions of austenite (Fig. 4(a)). This morphology was found in regions which had not been affected by weld thermal cycles due to subsequent weld pass. It constitutes major portion of the deposited weld metal. Ferrite of globular morphology was present in region near to the weld pass interface over which the subsequent weld pass was laid, in other
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Fig. 4. Microstructure of the as welded joint showing δ-ferrite with (a) vermicular and (b) globular morphologies.
Fig. 6. Variation of hardness along a weld pass interface. Fig. 5. Variation of hardness along X–X0 AND Y–Y0 sections.
words it can be said that this region is the ‘heat affected zone’ in the weld metal (Fig. 4(b)). The spread of this region was up to 700 μm from the weld pass interface. The variations of hardness across crown to root (Y–Y0 section in Fig. 3) and across base metal-heat affected zone to weld metal-heat affected zone-base metal (X–X0 section in Fig. 3) of the weld joint are shown in Fig. 5. Though hardness increased from crown to root of the weld metal (in the Y–Y0 section), there was appreciable scatter in the data. The weld metal in general exhibited higher hardness than the base metal due to its as-cast structure. The increase in hardness from crown to root of the joint can be attributed to the increased dislocation density generated by weld thermal cycles of the multipass welding. The appreciable scattering in hardness is considered due to the microstructural inhomogeneity across each weld pass. The hardness variation between successive weld passes is shown in Fig. 6. The globular δ-ferrite region which had formed as a result of thermal cycle produced by the subsequent weld pass, exhibited higher hardness than that of the vermicular ferrite region. There was similar variation in hardness of the weld metal across X–X0 section as in Y–Y0 . It is however important to note that the heat affected zone in the base metal exhibited higher hardness than the base metal, this again can be attributed to increased dislocation density produced by weld thermal cycle [11–13]. TEM micrograph of the vermicular and globular ferrite regions of the weld metal are shown in Fig. 7(a) and (b) respectively. The interconnected δ-ferrite, as observed from the optical micrograph (Fig. 3(a)), could be clearly seen along with free dislocations in both δ-ferrite and austenite in the vermicular ferrite region.
The substructure in the globular regions consisted of dislocation forest structure and isolated δ-ferrite with a globular morphology (Fig. 7(b)), giving rise to higher hardness (Fig. 6). As in case of the heat affected region of 316LN SS weld joints, the dislocation density was higher in this region subjected weld thermal cycle by the subsequent weld thermal cycle. Although through TEM analysis the spread of this region consisting of dislocation forest structure could not be established, hardness measurements indicated that this region matched with the region of globular ferrite morphology. 3.2. Creep rupture properties A comparison between creep rupture lives of the base metal and cross weld joint, creep tested at 923 K over at stresses in the range 120–225 MPa, is shown in Fig. 8. The rupture life of weld joints is substantially lower than that of the base metal, especially at relatively lower applied stresses. The difference in creep rupture strength between the base metal and weld joint enhanced with increase in rupture life. Failure of all the weld joints occurred in the weld metal region. The microstructural instability of δ-ferrite contained in weld region, as discussed subsequently, is considered for the reduction in creep rupture life of weld joint than the base metal. The variations of creep rupture ductility with creep rupture life of both the base metal and the weld joint are shown in Fig. 9. Creep ductility decreased with creep exposure for the weld joints whereas it increased for the base metal. The gradual reduction in creep rupture ductility in the weld joint can be attributed to the nucleation and growth of creep cavities associated with the
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Fig. 7. TEM micrograph of the as-weld material showing (a) vermicular interconnected δ-ferrite and (b) globular δ-ferrite with higher density of dislocations.
Fig. 10. Variation of residual δ-ferrite with rupture life. Fig. 8. Variation of rupture life with applied stress.
intermetallics. The base metal is more resistant to cavitation after longer creep exposure because the carbides in the grain boundaries coarsen and they offer very less resistance to grain boundary sliding. The residual δ-ferrite content steadily reduced with creep exposures and it could be ascertained that for the weld joint with longest rupture the transformation to brittle intermetallics is almost complete (Fig. 10). 3.3. Microstructural modification of weld metal on creep exposure
Fig. 9. Variation of rupture elongation with rupture life.
The elevated temperature transformation of δ-ferrite into carbides and intermetallics has been widely studied in austenitic stainless steel welds [14,15]. The transformation commences with the precipitation of M23C6 at the austenite/δ-ferrite interface. Prolonged exposure can lead to the precipitation of intermetallic phases such as s, χ and η. It has been argued that the amounts of C and N in the weld metal have definitive roles in the kinetics of precipitation of these intermetallics. While N inhibits the formation of all the aforesaid intermetallic phases, C on the other hand restricts the precipitation of s and η phases, but not the χ phase. This is because C is soluble in χ and the precipitation of this phase can happen simultaneously along with M23C6 [16]. Though precipitation of intermetallics can occur in austenite also, its kinetics is rather slow. The δ-ferrite has higher
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solubility for Cr, Mo which are the main constituents of the intermetallics. Further, the diffusion of these elements is higher in δ-ferrite when compared to austenite, because of its more open BCC structure. During creep exposure, the precipitation of the phases is enhanced than in thermal aging due to the stress induced diffusion of the elements [17]. In light of the above discussion it is clear that the essential factors for the precipitation of carbides are the diffusion of Cr and Mo. The formation of M23C6 depends on the diffusion of Cr, whereas diffusion of both Cr and Mo is essential for the precipitation of intermetallics [18]. It was also pointed out that continuous network of δ-ferrite results in higher δ-ferrite/austenite interface boundaries which act as nucleation sites for carbides and intermetallics. In the present investigation, the globular morphology consisted of isolated δ-ferrite which lowers interface boundaries. Thus the transformation of δ-ferrite into carbides and intermetallics is delayed in regions where delta ferrite exhibited globular morphology. Fig. 11 shows the variations of contents of Cr and Mo in the transformed phases with creep exposure, estimated by EDS attached with SEM, in both the globular and vermicular region of the weld metal. It can be clearly seen that the enrichment of Cr and Mo in the phases was higher in the vermicular regions when compared to the globular regions. Hence, it can be ascertained that the transformation of δ-ferrite into carbides and intermetallics could readily occur in the vermicular region than that in globular region. It should be also noted that multi-pass welding results in variations of the initial delta ferrite content in the globular and vermicular regions [9], and this can also influence the transformation kinetics.
3.4. Dislocation substructure evolution during creep The evolution of dislocation substructure was studied in detail for one short term test (applied stress 175 MPa, rupture life 1300 h) and one long term test (applied stress 120 MPa, rupture life 7820 h). The weld joint tested at 175 MPa showed dislocation pile ups and slip bands in the vermicular region (Fig. 12(a) and (b)). However in the globular region, subgrains could be seen (Fig. 12(c)). Precipitation of intermetallics was limited in this test condition. Fig. 13(a) shows the dislocation substructure in the vermicular region of the weld joint tested at 120 MPa. The cell structure which is a characteristic of recovery during creep of alloys with low stacking faults is clearly visible along with coarsened intermetallics. EDS spectra on these intermetallics revealed that they were Fe–Cr–Mo and Mo–Fe rich suggesting that they could be s and η respectively (Fig. 14). The TEM micrograph taken in the globular region also showed intermetallics and subgrains (Fig. 13(b)). However, the dislocation distribution inside the subgrains was higher in this region when compared to those in the vermicular region. The evolution of dislocation substructure during creep of austenitic steel welds is well documented [19,20]. At high stress levels, in a well annealed microstructure such as the one in the vermicular region of the weld metal in the current investigation – planar slip is the principal deformation mechanism and the matrix resistance to plastic deformation is quite negligible. A cold worked structure with higher dislocation density such as the one in the globular region, subgrains form which is the signature of recovery during creep deformation. Although prolonged creep exposures (creep exposure at lower stress levels) produced subgrains in the in both vermicular and globular regions, there were some remnant dislocations within the subgrains of globular regions. This suggests that the formation of subgrains is not complete in this region. The excessive size of the intermetallics would not have caused any significant interactions with the dislocations in both the regions [19]. Hence it can be assumed that the initial dislocation substructure is the decisive factor driving recovery in both the regions of the weld at this stress level. The initial dislocations density was less and there were free dislocations in the vermicular regions when compared to a more tangled dislocation structure in the globular region. The applied stress in this case was not high enough for recovery mechanisms such as dislocation climb to operate; hence formation of sub grains was delayed in the dislocation rich globular region. 3.5. Creep damage and failure.
Fig. 11. Variation of Cr and Mo in transformed regions with rupture life.
Microstructure of the weld joint tested at 175 MPa showed intragranular cracking in the vermicular regions near the weld
Fig. 12. (a) Dislocation pile up, (b) slip bands in vermicular region and (c) subgrains formation in globular region after creep testing at 175 MPa.
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Fig. 13. TEM micrographs of weld joint tested at 120 MPa (rupture life 7820 h) (a) vermicular region consisting subgrains and precipitate A with Fe–Cr–Mo and B with Fe–Mo enrichment and (b) globular region with subgrains.
Fig. 14. EDS spectra of (a) precipitate A showing enrichment of Fe–Cr–Mo and (b) precipitate B showing Fe–Mo enrichment.
Fig. 15. (a) Intergranular cracking occurred for at 175 MPa and (b) creep cavitation occurred along the intermetallic/matrix after creep testing at 120 MPa.
Fig. 16. (a) Fractograph of weld joint tested at 175 MPa, (b) region I showing interdendritic failure and (c) region D showing ductile failure.
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Fig. 17. (a) Fractograph of weld joint tested at 120 MPa, (b) region B showing cracks in brittle transformed intermetallics and (c) region D showing ductile failure.
pass interface (Fig. 15(a)) whereas regions with globular morphology did not exhibit any cracking. As discussed earlier there is a significant difference in the dislocation substructure in the two regions which is also reflected in the hardness values reported earlier. This results in stress concentration in the regions very near the interface, as a consequence there is tendency for intergranular cracking to occur in comparatively softer region, which in the present context refers to the vermicular region. Fig. 15(b) shows the SEM micrograph in the vermicular region of the weld joint tested at 120 MPa. Interconnected cavitation along the transformed regions is clearly visible in this region. As pointed out earlier, the morphology of δ-ferrite in these regions is favourable for transformation kinetics of in this region, and is more prone to creep cavitation. The origin of creep damage in the intermetallic/ matrix interface has been discussed by many researchers in austenitic stainless steel weld metal [21–23]. Figs. 16 and 17 show the fractographs weld joints tested at 175 and 120 MPa respectively. The aforementioned cracks which formed in the intergranular regions for the weld joint tested at 175 MPa, manifested as fibrous region (Fig. 16(b)). Whereas the weld joint tested at 120 MPa showed regions of extensive cracking in the transformed brittle intermetallics (Fig. 17(b)). Both the fractographs contained ductile features which could have possibly formed due to failure during ‘overload’ in the globular regions of the weld metal (Figs. 16 (c) and 17(c)). But it could be clearly seen that the extent of ductile region was more in the weld joint crept for 1300 h (applied stress 175 MPa) due to the incomplete transformation of δ-ferrite. 4. Conclusion Creep rupture properties of multi-pass SMAW 316LN SS was studied in the stress levels of 120–225 MPa at 923 K and the following conclusions have been drawn. 1. Depending on the creep testing conditions, the failure of the weld joints was attributed to two different phenomenons occurring due to the variations in morphology and dislocation substructure evolved due to multi-pass welding. 2. At higher stress levels the difference in dislocation substructure in the globular and vermicular regions led to stress concentration resulting in cracking of the softer vermicular region. 3. At low stress levels, creep cavitation was confined in the vermicular region due to the presence of transformed intermetallics and the strength inhomogeneity with the adjoining globular region. 4. Multi-pass welding produced regions of high dislocation density in the region consisting globular ferrite, rendering it more
resistant to creep deformation at high stress levels. At lower stress levels, where the failure was attributed to intermetallic creep cavitation, the transformation of δ-ferrite was delayed due to its globular morphology.
Acknowledgement The authors thank Dr. P.R. Vasudeva Rao, Director, Indira Gandhi Centre for Atomic Research, Kalpakkam and Dr. T. Jayakumar, Director, Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam, India for their continual support and guidance. References [1] S.L. Mannan, S.C. Chetal, Baldev Raj, S.B. Bhoje, Trans. Indian Inst. Met. 56 (2003) 155–178. [2] M.D. Mathew, K. Laha, V. Ganesan, Mater. Sci. Eng. A 535 (2012) 76–83. [3] V. Ganesan, M.D. Mathew, K. Bhanu Sankara Rao, Mater. Sci. Technol. 25 (2009) 614–618. [4] G.V. Prasad Reddy, R. Sandhya, K.B.S. Rao, S. Sankaran, Proc. Eng. 2 (2010) 2181–2188. [5] F.C. Hull, Weld. J. 60 (1967) 399–409. [6] J.J. Smith, R.A. Farrar, Int. Mater. Rev. 38 (1993) 25–51. [7] S.R. Keown, R.G. Thomas, Met. Sci. 15 (1981) 386–/392. [8] Y. Song, T.N Baker, N.A McPherson, Mater. Sci. Eng. A 212 (1996) 228–234. [9] C.A.P. Horton, J.K. Lai, Met. Sci. 14 (1980) 502–505. [10] Q. Auzoux, L. Allais, C. Caes, I. Monnet, A.F. Gourgues, A. Pineau, J. Nucl. Mater. 400 (2010) 127–137. [11] Naveena, V.D. Vijayanand, V. Ganesan, K. Laha, M.D. Mathew, Mater. Sci. Technol. (2013), http://dx.doi.org/10.1179/1743284713Y.0000000438. [12] P. Marshall, Austenitic Stainless Steels: Microstructure and Mechanical Properties, Elsevier Applied Science Publishers, London and New York, 1984. [13] C.F Etienne, D. Van Rossum, F. Roode, Proceedings of the International Conference on Engineering Aspects of Creep, September 1980, University of Sheffield, U.K., vol. 2, pp. 113–121. [14] K.H. Lo, C.H. Shek, J.K.I. Lai, Mater. Sci. Eng. R 65 (2009) 39–104. [15] G. Sasikala, M.D. Mathew, K.B.S. Rao, S.L. Mannan, Met. Mater. Trans. A 31 (2000) 1175–1185. [16] E. Folkhard, Welding Metallurgy of Stainless Steels, Springer Verlag, Wein, New York, 1986. [17] M.D. Mathew, G. Sasikala, T.P.S. Gill, S.L. Mannan, P. Rodriguez, Mater. Sci. Technol. 10 (1994) 1104–1108. [18] G. Sasikala, S.K. Ray, S.L. Mannan, Mater. Sci. Eng. A 359 (2003) 86–90. [19] M.D. Mathew, M. Sundararaman, S.L. Mannan, Mater. Trans. 38 (1997) 37–42. [20] J.R. Foulds, J. Monteff, Weld. J. 61 (1982) 189–196. [21] B.A. Senior, J. Mater. Sci. 25 (1990) 45–53. [22] G. Sasikala, S.K. Ray, S.L. Mannan, Acta Mater. 52 (2004) 5677–5686. [23] T. Shakhivel, M. Vasudevan, K. Laha, P. Parameswaran, K.S. Chandravathi, M.D. Mathew, A.K. Bhaduri, Mater. Sci. Eng. A 528 (2011) 6971–6980.