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Premature failure of superduplex stainless steel pipe by pitting in sea water environment
Engineering Failure Analysis 46 (2014) 134–139
Contents lists available at ScienceDirect
Engineering Failure Analysis journal homepage: www.elsevier.com/locate/engfailanal
Short communication
Premature failure of superduplex stainless steel pipe by pitting in sea water environment J. Smiderle a, J.M. Pardal b, S.S.M. Tavares b,⇑, A.C.N. Vidal c a
Universidade Federal do Rio de Janeiro, Departamento de Engenharia Metalúrgica e de Materiais, Rio de Janeiro, Brazil Universidade Federal Fluminense, Programa de Pós-Graduação em Engenharia Mecânica, Rua Passo da Pátria, 156, Niterói, RJ CEP 24210-240, Brazil c Pontifícia Universidade Católica do Rio de Janeiro, Instituto Tecnológico, Rio de Janeiro, Brazil b
a r t i c l e
i n f o
a b s t r a c t
Article history: Received 29 May 2014 Accepted 6 August 2014 Available online 19 August 2014
Ó 2014 Elsevier Ltd. All rights reserved.
Keywords: Superduplex stainless steels Pitting corrosion Sigma phase precipitation
1. Introduction Superduplex stainless steels (SDSS) allow high corrosion resistance and high mechanical strength. These two characteristics are the main reasons for the increasing use of this material in the oil and gas exploitation industry. Petrochemical industries, desalinization plants and modern oil and gas off-shore platforms were constructed with a large amount of facilities and equipment with SDSS [1–4]. It includes heat exchangers, pressure vessels, hydro-cyclones, tubbing, pipes and accessories. Corrosion resistant alloys (CRA’s) are frequently ranked by the pitting resistance equivalent (PRE), a composition based parameter given by [5]:
PRE ¼ %Cr þ 3:3ð%Mo þ 0:5ð%WÞÞ þ 16ð%NÞ Austenitic–ferritic steels with PRE > 40 are classified as superduplex. Steels with PRE < 40 are called duplex or lean duplex (without Mo). A direct correlation between the PRE and the critical pitting temperature (CPT) is currently presented [5]. However, microstructural features may drastically decrease the pitting resistance of duplex (DSS) and SDSS, despite of its PRE. Some examples of this are shown in previous works [5–9]. The typical microstructure of wrought SDSS consists of elongated islands of austenite and ferrite. It is well known that the best corrosion resistance and mechanical properties of duplex and superduplex stainless steels are obtained with about 50% of austenite (c) and ferrite (d). Deleterious phases, such as chromium nitrides (Cr2N) and intermetallics (r, R, v) must be avoided because they provoke severed decrease of corrosion resistance and mechanical properties. The temperatures intervals for precipitation of these phases have been extensively studied by many researchers [10–13]. ⇑ Corresponding author. Tel.: +55 21 2629 5584. E-mail address: [email protected] (S.S.M. Tavares). http://dx.doi.org/10.1016/j.engfailanal.2014.08.001 1350-6307/Ó 2014 Elsevier Ltd. All rights reserved.
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
135
Fig. 1. External side of the tube.
Fig. 2. Internal side of the tube.
This work deals with a failure analysis of superduplex pipe in a new platform for oil and gas transportation. The material failed prematurely, only 1 month after the platform start-up. Fig. 1 shows the 3.4 mm thickness tube failed with passant pits near a welded joint. The pits were concentrated in one of the tubes, named tube B in this work. It was also reported that the welded joint has been repaired. The environment conditions, seawater at room temperature, were not severe, considering the high pitting corrosion resistance (PRE > 40) of the SDSS. However, pitting corrosion was observed in one side of the joint, in a perimeter of about 120 mm of tube B, as shown in Fig. 2. The pits were nucleated in the inside wall. 2. Methodology The welded tube was cut for analysis. Chemical analysis by plasma spectroscopy was performed in the base and weld metals. Nitrogen of base metals was analyzed by combustion method with sparks. Fig. 3 shows in detail the through-thickness pit in tube B. A sample was carefully cut for metallographic analysis. The microstructure was investigated by optical microscopy, with samples prepared with electrolytic etching (3V, 20 s) in 10% KOH solution, or with Beraha’s etching (80 ml H2O, 20 ml HCl and 0.4 g of potassium metabisulfite). KOH etching is recommended to observe deleterious phases (r, v, R, . . .) in duplex and austenitic steels [14,15], while Beraha’s etching is used to quantify austenite and ferrite [1]. Scanning electron microscope (SEM) analysis was performed in specimens polished and not etched. Vickers hardness was measured with load of 5 kgf in both tubes and in the weld metal. The ferrite phase in tubes A and B, and in the weld metal was analyzed with a Helmut Fischer ferritoscope calibrated with standard duplex steel samples. The pitting corrosion resistance of tube B was analyzed by the measurement of CPT by the potentiostatic method (ASTM G-150) [16]. The tests were carried in a three electrode cell, with saturated calomel electrode as reference, Pt foil as counterelectrode and the welded joint as work electrode. Three regions were selected for polarization tests: weld metal (WM), tube A, and tube B. The cell potential was 0.7 VSCE. The temperature cell was raised with a rate of 1 °C/min, while the current density was recorded.
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J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
Fig. 3. Region of perforation by pitting in tube B showing specimen cut for analysis.
3. Results and discussion 3.1. Chemical analysis Table 1 shows the chemical compositions of tubes A and B and weld metal (root and cap pass). Nitrogen could not be analyzed by plasma spectroscopy in weld metal. Both tubes have chemical compositions typical of UNS S32760 superduplex stainless steel. However, the weld metal has Cu and W contents lower than 0.5%, which indicates that feed metals were ‘‘2509’’ wires without Cu and W. These wires are typically used to weld superduplex UNS S32750, which does not contain Cu and W in the nominal composition. 3.2. Microstructural analysis by LOM and SEM Fig.4 shows a profile of the welded joint with the pit located in tube B. The weld metal was divided in regions WM-R1, WM-R2 and WM-R3, as shown in Fig.4. The microstructures observed after Beraha’s etching are quite different in the three regions. WM-R3, probably the last pass, has a microstructure almost completely ferritic (Fig. 5(a)), WM-R1 has a balanced microstructure with 55% of austenite, (Fig. 5(b)), and WM-R2 has an intermediate microstructure (not shown), with about 40% of austenite. The microstructures of tubes A and B are shown in Fig. 6(a) and (b), respectively. Tube A has the typical microstructure of elongated islands of austenite and ferrite, produced by rolling and annealing or solution treatment. However tube B has a quite different microstructure in the region of pits. Besides ferrite and austenite, coarse globular or blocky intermetallic phases are observed. It can be also observed that the ferrite and austenite islands in tube B are finer than in tube A.
Table 1 Chemical compositions of tubes and weld metal. Tube
C
Cr
Ni
Mo
N
S
P
Cu
W
Tube A Tube B WM-root WM-cap
0.024 0.025 0.025 0.016
24.93 25.18 24.90 25.01
7.32 7.58 8.90 9.39
3.57 3.80 3.68 3.86
0.25 0.26 nd nd
0.004 0.004 0.006 0.004
0.022 0.022 0.019 0.018
0.59 0.57 0.32 0.17
0.51 0.50 0.11 0.06
"nd" means not determined.
Fig. 4. Profile of specimen cut for microstrucutural analysis.
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
137
Fig. 5. Microstructures of regions (a) WM-R3 and (b) WM-R1.
Fig. 6. Microstructures of tube (a) A and (b) B.
Fig. 7. Microstructure of tube B in detail.
Fig. 7 shows the microstructure of tube B with more detail. Ferrite and austenite were more attacked than the particles of intermetallic phase, in such a way that it was not possible to focus the three phases at the same time. So, the left side of Fig. 7 shows the austenite and ferrite in focus, while the right side shows the coarse precipitates in focus. SEM was used to identify the blocky particles by EDX analysis. This can be better achieved using the backscattered electrons (BSE) image, with specimens prepared by polishing and not etched [17]. Figs. 8(a and b) show images from a region of pitting. The chemical analysis of the particle of Fig. 8(b) and the ferrite and austenite phases are compared in Table 2. The chemical composition of the globular particles with higher Cr and Mo content is typical from r phase. Due to the higher Mo content, r is observed more brilliant than ferrite and austenite in the backscattered electrons image. As observed in Fig. 8(b), pit nucleation occurs near or around the blocky r particles. 3.3. Ferritoscope and microhardness Other indirect evidences of the r phase precipitation were obtained from the inspection with ferritoscope and hardness tests. Table 3 shows the results from tube A, weld metal and tube B close and far from the pitting region. Ferritoscope measures the ferrite number or ferrite fraction from the magnetic permeability of the sample. The decrease of ferrite indicates its
138
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
Fig. 8. SEM–BSE images from tube B, in a region of pitting corrosion. Specimens were polished preserving some pits.
Table 2 Semi-quantitative analysis by EDX of ferrite, austenite and coarse particles (r phase). (Average of three measurements; %Fe = bal.). Region
Cr
Mo
Ni
Si
Mn
Ferrite Austenite Globular particle (r)
26.8 24.9 30.3
4.9 3.4 9.2
6.1 7.4 4.8
0.8 0.4 2.0
0.5 0.7 0.2
Table 3 Ferritoscope and hardness measurements. Region
Ferritoscope reading (%d)
Hardness (HV)
Tube A Tube B (close to pits) Tube B (far from pitting area)
45.5 ± 3.2 20.4 ± 4.5 46.0 ± 2.7
275 ± 6 322 ± 19 282 ± 5
transformation into paramagnetic phases (r, c2, v). On the other hand, r phase precipitation provokes hardening only if present in amounts higher than about 10% [5]. It is worth noting that inspection of tube B far from the region of pitting did not indicate ferrite decrease. The microstructure at this region (not shown) was also free of r phase. However, due to its finer austenite and ferrite islands, the hardness was slightly higher than in tube A. In this case, the microstructure refinement also favors the kinetics of sigma phase precipitation in tube B. 3.4. Pitting corrosion test – CPT measurement As reported [5–9], r phase precipitation decrease the pitting resistance of austenitic and austenitic–ferritic steels. According to Nilsson [7], the critical pitting temperature (CPT) of a UNS S32750 decreases from 80 °C to less than 40 °C if
Tube B CPT = 32oC
Current density (A/cm 2)
0,0005 0,0004 0,0003 0,0002
32 oC
0,0001 0,0000 0
10
20
30
40
Temperature ( oC) Fig. 9. Curve for CPT measurement in tube B close to pits.
50
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
139
7–9% of r phase is formed. The amount of blocky r phase precipitated near the pits of tube B was not quantified, but the image from Fig. 7 and the hardness and ferritoscope results suggest that it was higher than 10% near the pitting areas. As shown in Fig. 9, the CPT measured in a sample from tube B near the pitting attack was 32 °C, confirming the previous results [7–9]. 4. Discussion The evidences show that tube B has undergone a localized overheating which caused intense r phase precipitation. Coarse blocky plates of r phase particles are produced in DSS and SDSS by exposure at temperatures as high as 900 °C [6,17,18]. Lower temperatures (700–900 °C) produce finer r particles, often associated with secondary austenite and v phase. d ? c2 + r eutectoid reaction is often proposed for r precipitation in austenitic–ferritic steels, but significant amounts of secondary austenite (c2) were not observed in tube B. Due to operational conditions, the localized overheating could not have happened under service. As reported, the welded joint had to be repaired. The weld metal does not contain r phase, but presents an heterogeneous microstructure, which indicates that it was not heat treated after welding. In view of these facts, the most likely explanation for the high temperature exposure is that a portion of tube B was heated to correct a previous distortion, before welding. Hot deformation to correct dimensional distortion is a common procedure in carbon and low alloy steels, but unacceptable for DSS and SDSS. The lack of knowledge of the staff responsible for the repair must have conducted to the strong mistake of heat the superduplex steel tube. 5. Conclusions and recommendation The pitting attack was caused by intense sigma phase precipitation in a portion of tube B. Tube A and the welded joint were free from r, but the microstructure of the weld metal was heterogeneous. The r phase precipitation was due to incorrect overheating of a portion of tube B, likely performed to deform and correct dimensional distortion before the repair welding. It was observed that r phase caused a detectable reduction of ferrite reading with ferritoscope. Field inspections can be conducted near the welded joints, mainly in those which were repaired. The portions with low ferrite content must be investigated with field metallography and/or non-destructive hardness tests. Tubes with detectable amounts of r phase must be repaired by substitution of the damaged portions, and welding according to standard specifications. Acknowledgements Authors acknowledge the Brazilian Research Agencies CAPES, CNPq, and FAPERJ for the financial support. References [1] Tavares SSM, Scandian C, Pardal JM, Luz TS, da Silva FJ. Failure analysis of duplex stainless steel tube weld used in flexible pipes in off shore oil production. Eng Fail Anal 2010;17:1500–6. [2] Lasebikan BA, Akisanya AR. Burst pressure of superduplex steel pipes subjected to combined axial tension, internal pressure and elevated temperature. Int J Press Ves Pip 2014;119:62–8. [3] Lee C-H, Chang K-H. Failure pressure of a pressurized girth-welded superduplex steel pipe in reverse osmosis desalination plants. Energy 2013;61(1):565–74. [4] Lynton VM, Laycock NJ, Thomsen SJ, Klumpers A. Eng Fail Anal 2004;11:243–56. [5] Gunn RN. Duplex stainless steels: microstructure, properties and applications. Cambridge: Abington Publishing; 2003. [6] Lopez N, Cid M, Puiggali. Influence of r phase on mechanical properties and corrosion resistance of duplex stainless steels. Corros Sci 1999;41:1615–31. [7] Nilsson JO. Influence of isothermal phase transformations on toughness and pitting corrosion of super duplex stainless steel SAF 2507. Mater Sci Technol 1993;9:545–54. [8] Kina AY, Tavares SSM, Souza VM, Lima LD, Corte RRA, Pardal JM. Influence of microstructure on the corrosion resistance of the duplex stainless steel UNS S31803. Mater Charact 2008;59:1127–32. [9] Potgieter JH. Influence of r phase on general and pitting corrosion resistance of SAF 2205 duplex stainless steel. Br Corros J 1992;27(3):219–23. [10] Chen TH, Weng KL, Yang JR. The effect of high-temperature exposure on the microstructural stability and toughness property in a 2205 duplex stainless steel. Mater Sci Eng 2002;A338:259–70. [11] Krull P, Pries H, Wohlfahrt H, Tösch J. Precipitation and toughness behavior of thick-walled weld joints made of duplex steel. Weld Cutting 1997;11:2–7. [12] Nilsson JO, Kangas P, Karlsson T, Wilson A. Mechanical Properties, microstructural stability and kinetics of r phase formation in 29Cr–6Ni–2Mo–0.38N superduplex stainless steel. Metall Mater Trans 2000;31A:35–45. [13] Nilsson JO. Overview super duplex stainless steels. Mater Sci Technol 1992;8:685–700. [14] Pardal JM, Carvalho SS, Barbosa C, Montenegro TR, Tavares SSM. Failure analysis of AISI 310S plate in an inert gas generator used in off-shore oil platform. Eng Fail Anal 2011;18:1435–44. [15] Pardal JM, Tavares SSM, Cindra Fonseca MP, Souza JA, Loureiro A, Moura EP. Modelling of deleterious phase precipitation during isothermal treatments in superduplex stainless steel. J Mater Sci 2010;45:616–23. [16] ASTM G-150-99. Standard test method for electrochemical critical pitting temperature testing of stainless steels. West Conshohocken, PA: ASTM International; 1999. [17] Pardal JM, Tavares SSM, Cindra Fonseca MP, Souza JÁ, Vieira LM, Abreu HFG. Deleterious phase precipitation on superduplex stainless steel UNS S32750 characterization by light optical and scanning electron microscopy. Mater Res 2010;13(3):401–7. [18] Pohl M, Storz O, Glogowski T. Effect of sigma phase morphology on the properties of duplex stainless steels. Microsc Microanal 2005;11(2):230–1.
Contents lists available at ScienceDirect
Engineering Failure Analysis journal homepage: www.elsevier.com/locate/engfailanal
Short communication
Premature failure of superduplex stainless steel pipe by pitting in sea water environment J. Smiderle a, J.M. Pardal b, S.S.M. Tavares b,⇑, A.C.N. Vidal c a
Universidade Federal do Rio de Janeiro, Departamento de Engenharia Metalúrgica e de Materiais, Rio de Janeiro, Brazil Universidade Federal Fluminense, Programa de Pós-Graduação em Engenharia Mecânica, Rua Passo da Pátria, 156, Niterói, RJ CEP 24210-240, Brazil c Pontifícia Universidade Católica do Rio de Janeiro, Instituto Tecnológico, Rio de Janeiro, Brazil b
a r t i c l e
i n f o
a b s t r a c t
Article history: Received 29 May 2014 Accepted 6 August 2014 Available online 19 August 2014
Ó 2014 Elsevier Ltd. All rights reserved.
Keywords: Superduplex stainless steels Pitting corrosion Sigma phase precipitation
1. Introduction Superduplex stainless steels (SDSS) allow high corrosion resistance and high mechanical strength. These two characteristics are the main reasons for the increasing use of this material in the oil and gas exploitation industry. Petrochemical industries, desalinization plants and modern oil and gas off-shore platforms were constructed with a large amount of facilities and equipment with SDSS [1–4]. It includes heat exchangers, pressure vessels, hydro-cyclones, tubbing, pipes and accessories. Corrosion resistant alloys (CRA’s) are frequently ranked by the pitting resistance equivalent (PRE), a composition based parameter given by [5]:
PRE ¼ %Cr þ 3:3ð%Mo þ 0:5ð%WÞÞ þ 16ð%NÞ Austenitic–ferritic steels with PRE > 40 are classified as superduplex. Steels with PRE < 40 are called duplex or lean duplex (without Mo). A direct correlation between the PRE and the critical pitting temperature (CPT) is currently presented [5]. However, microstructural features may drastically decrease the pitting resistance of duplex (DSS) and SDSS, despite of its PRE. Some examples of this are shown in previous works [5–9]. The typical microstructure of wrought SDSS consists of elongated islands of austenite and ferrite. It is well known that the best corrosion resistance and mechanical properties of duplex and superduplex stainless steels are obtained with about 50% of austenite (c) and ferrite (d). Deleterious phases, such as chromium nitrides (Cr2N) and intermetallics (r, R, v) must be avoided because they provoke severed decrease of corrosion resistance and mechanical properties. The temperatures intervals for precipitation of these phases have been extensively studied by many researchers [10–13]. ⇑ Corresponding author. Tel.: +55 21 2629 5584. E-mail address: [email protected] (S.S.M. Tavares). http://dx.doi.org/10.1016/j.engfailanal.2014.08.001 1350-6307/Ó 2014 Elsevier Ltd. All rights reserved.
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
135
Fig. 1. External side of the tube.
Fig. 2. Internal side of the tube.
This work deals with a failure analysis of superduplex pipe in a new platform for oil and gas transportation. The material failed prematurely, only 1 month after the platform start-up. Fig. 1 shows the 3.4 mm thickness tube failed with passant pits near a welded joint. The pits were concentrated in one of the tubes, named tube B in this work. It was also reported that the welded joint has been repaired. The environment conditions, seawater at room temperature, were not severe, considering the high pitting corrosion resistance (PRE > 40) of the SDSS. However, pitting corrosion was observed in one side of the joint, in a perimeter of about 120 mm of tube B, as shown in Fig. 2. The pits were nucleated in the inside wall. 2. Methodology The welded tube was cut for analysis. Chemical analysis by plasma spectroscopy was performed in the base and weld metals. Nitrogen of base metals was analyzed by combustion method with sparks. Fig. 3 shows in detail the through-thickness pit in tube B. A sample was carefully cut for metallographic analysis. The microstructure was investigated by optical microscopy, with samples prepared with electrolytic etching (3V, 20 s) in 10% KOH solution, or with Beraha’s etching (80 ml H2O, 20 ml HCl and 0.4 g of potassium metabisulfite). KOH etching is recommended to observe deleterious phases (r, v, R, . . .) in duplex and austenitic steels [14,15], while Beraha’s etching is used to quantify austenite and ferrite [1]. Scanning electron microscope (SEM) analysis was performed in specimens polished and not etched. Vickers hardness was measured with load of 5 kgf in both tubes and in the weld metal. The ferrite phase in tubes A and B, and in the weld metal was analyzed with a Helmut Fischer ferritoscope calibrated with standard duplex steel samples. The pitting corrosion resistance of tube B was analyzed by the measurement of CPT by the potentiostatic method (ASTM G-150) [16]. The tests were carried in a three electrode cell, with saturated calomel electrode as reference, Pt foil as counterelectrode and the welded joint as work electrode. Three regions were selected for polarization tests: weld metal (WM), tube A, and tube B. The cell potential was 0.7 VSCE. The temperature cell was raised with a rate of 1 °C/min, while the current density was recorded.
136
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
Fig. 3. Region of perforation by pitting in tube B showing specimen cut for analysis.
3. Results and discussion 3.1. Chemical analysis Table 1 shows the chemical compositions of tubes A and B and weld metal (root and cap pass). Nitrogen could not be analyzed by plasma spectroscopy in weld metal. Both tubes have chemical compositions typical of UNS S32760 superduplex stainless steel. However, the weld metal has Cu and W contents lower than 0.5%, which indicates that feed metals were ‘‘2509’’ wires without Cu and W. These wires are typically used to weld superduplex UNS S32750, which does not contain Cu and W in the nominal composition. 3.2. Microstructural analysis by LOM and SEM Fig.4 shows a profile of the welded joint with the pit located in tube B. The weld metal was divided in regions WM-R1, WM-R2 and WM-R3, as shown in Fig.4. The microstructures observed after Beraha’s etching are quite different in the three regions. WM-R3, probably the last pass, has a microstructure almost completely ferritic (Fig. 5(a)), WM-R1 has a balanced microstructure with 55% of austenite, (Fig. 5(b)), and WM-R2 has an intermediate microstructure (not shown), with about 40% of austenite. The microstructures of tubes A and B are shown in Fig. 6(a) and (b), respectively. Tube A has the typical microstructure of elongated islands of austenite and ferrite, produced by rolling and annealing or solution treatment. However tube B has a quite different microstructure in the region of pits. Besides ferrite and austenite, coarse globular or blocky intermetallic phases are observed. It can be also observed that the ferrite and austenite islands in tube B are finer than in tube A.
Table 1 Chemical compositions of tubes and weld metal. Tube
C
Cr
Ni
Mo
N
S
P
Cu
W
Tube A Tube B WM-root WM-cap
0.024 0.025 0.025 0.016
24.93 25.18 24.90 25.01
7.32 7.58 8.90 9.39
3.57 3.80 3.68 3.86
0.25 0.26 nd nd
0.004 0.004 0.006 0.004
0.022 0.022 0.019 0.018
0.59 0.57 0.32 0.17
0.51 0.50 0.11 0.06
"nd" means not determined.
Fig. 4. Profile of specimen cut for microstrucutural analysis.
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
137
Fig. 5. Microstructures of regions (a) WM-R3 and (b) WM-R1.
Fig. 6. Microstructures of tube (a) A and (b) B.
Fig. 7. Microstructure of tube B in detail.
Fig. 7 shows the microstructure of tube B with more detail. Ferrite and austenite were more attacked than the particles of intermetallic phase, in such a way that it was not possible to focus the three phases at the same time. So, the left side of Fig. 7 shows the austenite and ferrite in focus, while the right side shows the coarse precipitates in focus. SEM was used to identify the blocky particles by EDX analysis. This can be better achieved using the backscattered electrons (BSE) image, with specimens prepared by polishing and not etched [17]. Figs. 8(a and b) show images from a region of pitting. The chemical analysis of the particle of Fig. 8(b) and the ferrite and austenite phases are compared in Table 2. The chemical composition of the globular particles with higher Cr and Mo content is typical from r phase. Due to the higher Mo content, r is observed more brilliant than ferrite and austenite in the backscattered electrons image. As observed in Fig. 8(b), pit nucleation occurs near or around the blocky r particles. 3.3. Ferritoscope and microhardness Other indirect evidences of the r phase precipitation were obtained from the inspection with ferritoscope and hardness tests. Table 3 shows the results from tube A, weld metal and tube B close and far from the pitting region. Ferritoscope measures the ferrite number or ferrite fraction from the magnetic permeability of the sample. The decrease of ferrite indicates its
138
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
Fig. 8. SEM–BSE images from tube B, in a region of pitting corrosion. Specimens were polished preserving some pits.
Table 2 Semi-quantitative analysis by EDX of ferrite, austenite and coarse particles (r phase). (Average of three measurements; %Fe = bal.). Region
Cr
Mo
Ni
Si
Mn
Ferrite Austenite Globular particle (r)
26.8 24.9 30.3
4.9 3.4 9.2
6.1 7.4 4.8
0.8 0.4 2.0
0.5 0.7 0.2
Table 3 Ferritoscope and hardness measurements. Region
Ferritoscope reading (%d)
Hardness (HV)
Tube A Tube B (close to pits) Tube B (far from pitting area)
45.5 ± 3.2 20.4 ± 4.5 46.0 ± 2.7
275 ± 6 322 ± 19 282 ± 5
transformation into paramagnetic phases (r, c2, v). On the other hand, r phase precipitation provokes hardening only if present in amounts higher than about 10% [5]. It is worth noting that inspection of tube B far from the region of pitting did not indicate ferrite decrease. The microstructure at this region (not shown) was also free of r phase. However, due to its finer austenite and ferrite islands, the hardness was slightly higher than in tube A. In this case, the microstructure refinement also favors the kinetics of sigma phase precipitation in tube B. 3.4. Pitting corrosion test – CPT measurement As reported [5–9], r phase precipitation decrease the pitting resistance of austenitic and austenitic–ferritic steels. According to Nilsson [7], the critical pitting temperature (CPT) of a UNS S32750 decreases from 80 °C to less than 40 °C if
Tube B CPT = 32oC
Current density (A/cm 2)
0,0005 0,0004 0,0003 0,0002
32 oC
0,0001 0,0000 0
10
20
30
40
Temperature ( oC) Fig. 9. Curve for CPT measurement in tube B close to pits.
50
J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139
139
7–9% of r phase is formed. The amount of blocky r phase precipitated near the pits of tube B was not quantified, but the image from Fig. 7 and the hardness and ferritoscope results suggest that it was higher than 10% near the pitting areas. As shown in Fig. 9, the CPT measured in a sample from tube B near the pitting attack was 32 °C, confirming the previous results [7–9]. 4. Discussion The evidences show that tube B has undergone a localized overheating which caused intense r phase precipitation. Coarse blocky plates of r phase particles are produced in DSS and SDSS by exposure at temperatures as high as 900 °C [6,17,18]. Lower temperatures (700–900 °C) produce finer r particles, often associated with secondary austenite and v phase. d ? c2 + r eutectoid reaction is often proposed for r precipitation in austenitic–ferritic steels, but significant amounts of secondary austenite (c2) were not observed in tube B. Due to operational conditions, the localized overheating could not have happened under service. As reported, the welded joint had to be repaired. The weld metal does not contain r phase, but presents an heterogeneous microstructure, which indicates that it was not heat treated after welding. In view of these facts, the most likely explanation for the high temperature exposure is that a portion of tube B was heated to correct a previous distortion, before welding. Hot deformation to correct dimensional distortion is a common procedure in carbon and low alloy steels, but unacceptable for DSS and SDSS. The lack of knowledge of the staff responsible for the repair must have conducted to the strong mistake of heat the superduplex steel tube. 5. Conclusions and recommendation The pitting attack was caused by intense sigma phase precipitation in a portion of tube B. Tube A and the welded joint were free from r, but the microstructure of the weld metal was heterogeneous. The r phase precipitation was due to incorrect overheating of a portion of tube B, likely performed to deform and correct dimensional distortion before the repair welding. It was observed that r phase caused a detectable reduction of ferrite reading with ferritoscope. Field inspections can be conducted near the welded joints, mainly in those which were repaired. The portions with low ferrite content must be investigated with field metallography and/or non-destructive hardness tests. Tubes with detectable amounts of r phase must be repaired by substitution of the damaged portions, and welding according to standard specifications. Acknowledgements Authors acknowledge the Brazilian Research Agencies CAPES, CNPq, and FAPERJ for the financial support. References [1] Tavares SSM, Scandian C, Pardal JM, Luz TS, da Silva FJ. Failure analysis of duplex stainless steel tube weld used in flexible pipes in off shore oil production. Eng Fail Anal 2010;17:1500–6. [2] Lasebikan BA, Akisanya AR. 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