Engineering Failure Analysis 46 (2014) 134–139

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Engineering Failure Analysis

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Short communication

Premature failure of superduplex stainless steel pipe by pittingin sea water environment

http://dx.doi.org/10.1016/j.engfailanal.2014.08.0011350-6307/� 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding author. Tel.: +55 21 2629 5584.E-mail address: ssmtavares@terra.com.br (S.S.M. Tavares).

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, Brazilb Universidade Federal Fluminense, Programa de Pós-Graduação em Engenharia Mecânica, Rua Passo da Pátria, 156, Niterói, RJ CEP 24210-240, Brazilc Pontifícia Universidade Católica do Rio de Janeiro, Instituto Tecnológico, Rio de Janeiro, Brazil

a r t i c l e i n f o

Article history:Received 29 May 2014Accepted 6 August 2014Available online 19 August 2014

Keywords:Superduplex stainless steelsPitting corrosionSigma phase precipitation

a b s t r a c t

� 2014 Elsevier Ltd. All rights reserved.

1. Introduction

Superduplex stainless steels (SDSS) allow high corrosion resistance and high mechanical strength. These two character-istics are the main reasons for the increasing use of this material in the oil and gas exploitation industry. Petrochemicalindustries, desalinization plants and modern oil and gas off-shore platforms were constructed with a large amount of facil-ities and equipment with SDSS [1–4]. It includes heat exchangers, pressure vessels, hydro-cyclones, tubbing, pipes andaccessories.

Corrosion resistant alloys (CRA’s) are frequently ranked by the pitting resistance equivalent (PRE), a composition basedparameter given by [5]:

PRE ¼ %Crþ 3:3ð%Moþ 0:5ð%WÞÞ þ 16 %Nð Þ

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. Someexamples 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 thebest 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 beavoided because they provoke severed decrease of corrosion resistance and mechanical properties. The temperatures inter-vals for precipitation of these phases have been extensively studied by many researchers [10–13].

Fig. 1. External side of the tube.

Fig. 2. Internal side of the tube.

J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139 135

This work deals with a failure analysis of superduplex pipe in a new platform for oil and gas transportation. The materialfailed prematurely, only 1 month after the platform start-up. Fig. 1 shows the 3.4 mm thickness tube failed with passant pitsnear a welded joint. The pits were concentrated in one of the tubes, named tube B in this work. It was also reported that thewelded joint has been repaired.

The environment conditions, seawater at room temperature, were not severe, considering the high pitting corrosion resis-tance (PRE > 40) of the SDSS. However, pitting corrosion was observed in one side of the joint, in a perimeter of about120 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 weldmetals. Nitrogen of base metals was analyzed by combustion method with sparks. Fig. 3 shows in detail the through-thick-ness pit in tube B. A sample was carefully cut for metallographic analysis. The microstructure was investigated by opticalmicroscopy, 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]. Scan-ning 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 (ASTMG-150) [16]. The tests were carried in a three electrode cell, with saturated calomel electrode as reference, Pt foil as counter-electrode and the welded joint as work electrode. Three regions were selected for polarization tests: weld metal (WM), tubeA, and tube B. The cell potential was 0.7 VSCE. The temperature cell was raised with a rate of 1 �C/min, while the currentdensity was recorded.

Fig. 3. Region of perforation by pitting in tube B showing specimen cut for analysis.

136 J. Smiderle et al. / Engineering Failure Analysis 46 (2014) 134–139

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 beanalyzed by plasma spectroscopy in weld metal. Both tubes have chemical compositions typical of UNS S32760 superduplexstainless 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 containCu 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 threeregions. WM-R3, probably the last pass, has a microstructure almost completely ferritic (Fig. 5(a)), WM-R1 has a balancedmicrostructure with 55% of austenite, (Fig. 5(b)), and WM-R2 has an intermediate microstructure (not shown), with about40% of austenite.

The microstructures of tubes A and B are shown in Fig. 6(a) and (b), respectively. Tube A has the typical microstructure ofelongated islands of austenite and ferrite, produced by rolling and annealing or solution treatment. However tube B has aquite different microstructure in the region of pits. Besides ferrite and austenite, coarse globular or blocky intermetallicphases are observed. It can be also observed that the ferrite and austenite islands in tube B are finer than in tube A.

Table 1Chemical compositions of tubes and weld metal.

Tube C Cr Ni Mo N S P Cu W

Tube A 0.024 24.93 7.32 3.57 0.25 0.004 0.022 0.59 0.51Tube B 0.025 25.18 7.58 3.80 0.26 0.004 0.022 0.57 0.50WM-root 0.025 24.90 8.90 3.68 nd 0.006 0.019 0.32 0.11WM-cap 0.016 25.01 9.39 3.86 nd 0.004 0.018 0.17 0.06

"nd" means not determined.

Fig. 4. Profile of specimen cut for microstrucutural analysis.

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.

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Fig. 7 shows the microstructure of tube B with more detail. Ferrite and austenite were more attacked than the particles ofintermetallic 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. 7shows 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 elec-trons (BSE) image, with specimens prepared by polishing and not etched [17]. Figs. 8(a and b) show images from a region ofpitting. The chemical analysis of the particle of Fig. 8(b) and the ferrite and austenite phases are compared in Table 2. Thechemical composition of the globular particles with higher Cr and Mo content is typical from r phase. Due to the higher Mocontent, 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 hardnesstests. Table 3 shows the results from tube A, weld metal and tube B close and far from the pitting region. Ferritoscope mea-sures the ferrite number or ferrite fraction from the magnetic permeability of the sample. The decrease of ferrite indicates its

Fig. 8. SEM–BSE images from tube B, in a region of pitting corrosion. Specimens were polished preserving some pits.

Table 2Semi-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 26.8 4.9 6.1 0.8 0.5Austenite 24.9 3.4 7.4 0.4 0.7Globular particle (r) 30.3 9.2 4.8 2.0 0.2

Table 3Ferritoscope and hardness measurements.

Region Ferritoscope reading(%d)

Hardness(HV)

Tube A 45.5 ± 3.2 275 ± 6Tube B (close to pits) 20.4 ± 4.5 322 ± 19Tube B (far from pitting area) 46.0 ± 2.7 282 ± 5

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transformation into paramagnetic phases (r, c2, v). On the other hand, r phase precipitation provokes hardening only ifpresent in amounts higher than about 10% [5]. It is worth noting that inspection of tube B far from the region of pittingdid not indicate ferrite decrease. The microstructure at this region (not shown) was also free of r phase. However, due toits finer austenite and ferrite islands, the hardness was slightly higher than in tube A. In this case, the microstructure refine-ment 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

0 10 20 30 40 500,0000

0,0001

0,0002

0,0003

0,0004

0,0005

Temperature (oC)

Cur

rent

den

sity

(A/c

m2 )

Tube BCPT = 32oC

32oC

Fig. 9. Curve for CPT measurement in tube B close to pits.

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 theimage from Fig. 7 and the hardness and ferritoscope results suggest that it was higher than 10% near the pitting areas. Asshown 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 vphase. d ? c2 + r eutectoid reaction is often proposed for r precipitation in austenitic–ferritic steels, but significantamounts 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 weldedjoint had to be repaired. The weld metal does not contain r phase, but presents an heterogeneous microstructure, whichindicates that it was not heat treated after welding. In view of these facts, the most likely explanation for the high temper-ature exposure is that a portion of tube B was heated to correct a previous distortion, before welding. Hot deformation tocorrect 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 super-duplex 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 jointwere 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 correctdimensional distortion before the repair welding.

It was observed that r phase caused a detectable reduction of ferrite reading with ferritoscope. Field inspections can beconducted near the welded joints, mainly in those which were repaired. The portions with low ferrite content must be inves-tigated with field metallography and/or non-destructive hardness tests. Tubes with detectable amounts of r phase must berepaired 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.

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