Magnetism variance analysis of Super304H austenitic stainless steel after service

Magnetism variance analysis of Super304H austenitic stainless steel after service

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Engineering Failure Analysis xxx (xxxx) xxx–xxx

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Magnetism variance analysis of Super304H austenitic stainless steel after service Wei Wanga, Xiang Huangb, Wen Sheng Lia, Yan Gaob,⁎ a b

Electric Power Research Institute of Guangdong Power Grid Co. Ltd., 510800, PR China School of Materials Science and Engineering, South China University of Technology, Guangzhou 510641, PR China



Keywords: Super304H steel High temperature service Grain size Cr-depleted zone Martensite

Magnetism variance of Super304H austenitic stainless steel superheater tubes after 36891h and 41000h service in the ultra-supercritical boilers was investigated by using optical microscope, scanning electron microscope, X-ray diffractometer, PPMS system (vibrating sample magnetometer, VSM), EBSD and electrochemical workstation. The results show that grain size of the 36891h serviced tube is extremely large with grade 2, and magnetization was detected in the tube, which was confirmed with a small amount of ferromagnetic phase distributing along some grain boundaries by EBSD observation, while the tube of 41000h service has a small grain size of grade 7–8 and no ferromagnetic phase or magnetization is present. The morphologies after DLEPR test show that the intergranular corrosion susceptibility of the 41000h tube is higher, but the depth and width of the Cr-depleted zones at grain boundaries are significantly smaller than the 36981h tube. The precipitates in Super304H after service did not cause significant changes in its magnetic properties. Coarse grains, which led to reduction of the diffusion channels, made Cr difficult to diffuse from the matrix to Cr-depleted zones. As a result, the Cr content in the Crdepleted zones maintained a lower level and resulted in the increase of Ms of the Cr-depleted zones, which induced martensitic transformation of Super304H tubes during the shut-down to room temperature of the boiler.

1. Introduction With the great increase of steam parameters (temperature and pressure) for the ultra-supercritical boilers, traditional TP304H, TP347H austenitic heat resistant steels are unable to meet the demand for high temperature properties when serving at 600 °C. Therefore, a novel Super304H austenitic heat-resistant steel, based on TP304H steel, has been developed by Japan Sumitomo Metal Company and Mitsubishi Heavy Industries Ltd. Due to its excellent high temperature mechanical properties, microstructure stability and excellent high temperature oxidation resistance, Super304H steel has been widely applied in ultra-supercritical units for superheater and reheater tubes whose wall temperature is around 600–650 °C and has shown good overall performance and achieved good economic benefits [1–3]. So far, over 7000 t Super304H steel has been used and more than 100 ultra-supercritical units in China in operation or under construction have adopt Super304H steel as the material of superheater and reheater tubes. It is well known that severe oxidation will inevitably happen when austenitic stainless steel serves in fossil fuel power plants under high temperature. Magnetic flux is often adopted to characterize the surface oxides and the pile up of oxides of austenitic heat resistant steel tubes used in fossil fuel power plants. Austenitic stainless steel is typical weak magnetic material and its magnetic

Corresponding author. E-mail address: [email protected] (Y. Gao). Received 11 July 2016; Accepted 8 May 2017 1350-6307/ © 2017 Published by Elsevier Ltd.

Please cite this article as: Huang, X., Engineering Failure Analysis (2017),

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permeability is very close to air permeability, while the oxides is ferromagnetic. Therefore, the thickness of the inner wall oxides and oxide pile up caused by oxide exfoliation can be determined according to magnetization intensity [4]. In an ultra-supercritical boiler A in China, the superheater tubes of Super304H steel which had run for 36891h, was detected by magnetic flux inspection to show obvious magnetic characterization, while no obvious magnetization had been detected in the inspection two years ago. Anatomical examination of the tube showed that the thickness of oxides was in a normal level and no oxide pile up was observed. However, in another ultra-supercritical boiler B which had serviced for 41000h, Super304H steel used for superheater tubes showed no obvious magnetization and the thickness of its oxide film was close to the tube serviced for 36891h in boiler A. The different change of magnetic flux in Super304H austenitic heat-resistant steels during high temperature service is worth investigating. When an austenitic heat-resistant steel runs at high temperature, oxide film and deformation may cause magnetization, meanwhile precipitations and martensitic transformation may also lead to magnetic flux change. Studies have shown that in the process of rapid cooling or cold deformation, heat-resistant austenitic steel with paramagnetic face-centered cubic structure may induce martensitic transformation to form a non-ferromagnetic “ε” phase (hexagonal) or ferromagnetic “α” phase (body-centered cubic) [5–6]. Since the as-received Super304H tubes showed no obvious magnetic flux, it is generally believed that three reasons may lead to the magnetization without deformation [7–8]: (1) Precipitations during the service process; (2) increase of Ms, caused by Cr depletion near grain boundaries, leading to martensitic transformation and the formation of ferromagnetic martensite; (3) thick oxide film or oxide pile-up. What causes the change of magnetic flux is a quite complicated issue. Therefore, the magnetism variance of Super304H austenitic stainless steel after service was investigated in this paper to find out what causes the magnetization change and distinct magnetic states of the two tubes whose service environment is similar, with serve time of 36891h in unit A and 41000h in unit B. 2. Material and methods Super304H steel tube serviced for 36891h with high magnetic flux and the tube serviced for 41000h with low magnetic flux were analyzed. Chemical composition of the tubes is shown in Table 1. Service parameters of both tubes are similar: designed steam outlet pressure and temperature of the tube serviced for 36891h is 26.15 MPa & 605 °C and is 26.25 MPa & 600 °C for the tube serviced for 41000h. Specification of the tubes serviced for 36891h and 41000h is ϕ51 mm × 9.5 mm and ϕ45 mm × 9 mm respectively. After removing the oxide layer on the tubes, magnetization intensity and coercivity of the tubes were measured to characterize the magnetization by using VSM (vibrating sample magnetometer). Microstructures of the specimens were observed by using optical microscope and FESEM Zeiss Supra-40 scanning electron microscope (SEM). Statistic volume fraction of Nb-rich phase was carried out by using Image-Pro-Plus software. Electron backscatter diffraction (EBSD) technique [9] was used to verify whether ferromagnetic “α” phase existed in the service tubes using. Preparation of the specimens for EBSD observation was as follows: after mechanical polishing, electrolytic polishing was carried out in solution (HClO4:C2H5OH = 1:9) at 0 °C for 30 s with a potential of 22 V and a current of 1.5–2.0 A. EBSD observation was made by using the Oxford EBSD system on NOVA NANOSEM 430. Cr depletion at grain boundary directly affects intergranular corrosion susceptibility (IGCS) which can be characterized by means of double loop electrochemical potential reactivation (DL-EPR) method [10]. The ratio of maximum reactivation current density (Ir) in reverse scan over the maximum activation current density (Ia) in forward scan in the DL-EPR curve is used to characterize the IGCS of materials, abbreviated as DOS (Degree of Sensitization). The greater the DOS value, the higher the susceptibility to intergranular corrosion. DL-EPR tests were conducted using a potentiostat workstation PGSTAT30. A saturated calomel electrode (SCE) and a graphite electrode were used as the reference and counter electrodes, respectively. The specimens, acting as working electrodes, were embedded in epoxy resin with an exposure area of 10 × 10 mm2 and ground using emery papers to 1200 grit. During DL-EPR tests, the working electrode was immersed in a 0.5 M H2SO4 + 0.01 M KSCN solution at 35 °C in a water bath to reach the steady-state open circuit potential (OCP) and then anodically polarized from − 500 mVSCE to +300 mVSCE (in the passive region). Reverse scan was made in the same voltage range to the OCP at a scan rate of 1.67 mV/s. After DL-EPR tests, the specimens were ultrasonic cleaned and corrosion morphologies were observed by SEM. 3. Results and discussion Since no oxide pile-up was observed during macro inspection for both the 36891h and 41000h serviced tubes with similar oxide film thickness and no magnetic flux was detected in the 41000h tube, the possibility caused by oxide pile up of the magnetic flux detected in the 36891h tube was debarred. Therefore two possibilities remain which may cause the magnetization: precipitations during service and increase of Ms caused by Cr depletion near grain boundaries. Table 1 Chemical composition of serviced Super304H tubes (mass fraction, %). No.













36891h service 41000h service

0.077 0.082

0.28 0.28

0.89 0.94

0.021 0.020

0.001 0.004

18.45 18.65

8.71 8.55

0.51 0.57

2.85 2.93

– –

– –

0.018 0.005


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Fig. 1. Dependence of magnetization on magnetic field intensity of Super304H steels a) as-received; (b) 36891h; (c) 41000h.

3.1. Magnetization of serviced Super304H steel To quantify the magnetism variance between Super304H tube serviced for 36891h and other tubes, the magnetic hysteresis loops of as-received and serviced tubes were measured using VSM and are shown in Fig. 1. When the linear portion of the magnetization curve is extrapolated to zero magnetic field intensity, the corresponding magnetization intensity is the specific saturated magnetization intensity (M). After magnetization of materials is saturated, the magnetization will not return to zero when external magnetic field returns to zero. Only when an opposite magnetic field against the original magnetic field is applied can the 3

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Table 2 Magnetic properties of Super304H tubes.

Magnetization intensity/emu/g Coercive force/Oe




0.09 145

0.20 95

0.75 45

magnetization returns to zero. The intensity of the opposite magnetic field is called coercive force. The saturated magnetization intensity and coercive force of the tubes are shown in Table 2. The magnetization intensity of the 41000h tube is 0.1–0.2 emu/g, which is close to that of the as-received tube with single austenite microstructure while the magnetization intensity of the tube serviced for 36891h reaches 0.75 emu/g and its coercive force is apparently smaller than the tube serviced for 41000h. The tube serviced for 36891h can be easily magnetized, meaning that its single austenite microstructure had changed during service and might have led to magnetic transformation. 3.2. Microstructure 3.2.1. Grain size The optical microstructures of the serviced specimens etched by the aqua regia are shown in Fig. 2. Their microstructure is all composed of γ phase and a small amount of bulk primary Nb-rich phase, with a big difference that the specimen of 41000h has a grain size grade of 7–8 which is close to that of the as-received tubes, while the specimen of 36891h has a much larger grain size of grade 2 (grain diameter of about 200 μm). 3.2.2. Precipitations a) Primary bulk Nb-rich phase Addition of Nb in Super304H aims at forming dispersive Nb(C, N) precipitates to strengthen the matrix and decrease the solute C content in the matrix which may lead to intergranular corrosion susceptibility. However, primary bulk Nb-rich precipitates in Super304H steel make Nb not able to precipitate uniformly and dispersively with C and N during solid solution treatment, which is bad for both strengthening effect and C constraint effect. Primary bulk Nb-rich phase in polished state was observed by SEM as shown in Fig. 3. Bulk primary Nb-rich phase exists in both serviced tubes and its volume fraction in 36891h and 41000h tubes, analyzed by Image-Pro-Plus software, is 0.29% and 0.31% respectively, which are close to each other. b) M23C6 M23C6 is one of the main precipitates in Super304H steel during service. Precipitation of M23C6 at grain boundaries will not only affect the high temperature properties, but also increase the intergranular corrosion susceptibility of Super304H steel. The SEM morphologies of the Super304H specimens serviced for 36891h and 41000h, are shown in Fig. 4. M23C6 carbides precipitating along grain boundaries in the tube of 36891h get coarsened and aggregated apparently compared to the tube of 41000h. It is worth noting that corrosion grooves along grain boundaries can be observed in Fig. 4(a) of the specimen serviced for 36891h while such morphology is not observed in the specimen serviced for 41000h. 3.3. DL-EPR test and corrosion morphology Compared to the TP304H steel, Super304H steel has higher C content, which favors more M23C6 formation at grain boundaries

Fig. 2. Metallographic morphologies of Super304H steels serviced for (a) 36891h; (b) 41000h.


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Fig. 3. SEM morphologies of polished Super304H steels serviced for (a) 36891h; (b) 41000h.

and thus more Cr depletion zones near grain boundaries and consequently the increase of intergranular corrosion susceptibility. According to the DL-EPR curves of Super304H steel shown in Fig. 5, the DOS of the specimen serviced for 36891h is 28.2%, much lower than the DOS of 36.1% of the specimen serviced for 41000h, indicating less carbides and less Cr depletion zones in the 36891h tube. When the corrosion morphologies after DL-EPR tests are observed, as shown in Fig. 6, however, the width of the Cr-depleted grooves at grain boundaries of the 36891h tube is found to be obviously larger than the 41000h tube, which can be interpreted that the coarse grains of the 36891h tube led to reduction of diffusion channels and made Cr difficult to diffuse from the matrix to the Crdepleted zones. As a result, the Cr content near the grain boundaries maintained a lower level and led to wider and deeper corrosion grooves near grain boundaries in the 36891h tube.

4. Discussion 4.1. Effect of precipitates on magnetism variance of Super304H steel Nb in Super304H plays the roles of strengthening the matrix by forming dispersive Nb(C, N) during solution treatment and by constraining M23C6 formation at the same time. More bulk primary Nb-rich phase will lead to less dispersive fine Nb(C, N) and more M23C6 precipitation. The tubes serviced for 36891h and 41000h have similar volume fraction of primary bulk Nb-rich phase, therefore primary bulk Nb-rich phase should not be the cause of magnetic difference. Study shows that magnetic transition did not occur even when the M23C6 precipitation had reached a saturation in the Super304H tube serviced for longer time [11] which means that M23C6 carbides themselves will not lead to magnetic transition. However, Cr depletion zones near M23C6 carbides, especially when M23C6 carbides get coarsened, will corrode preferentially. Wider corrosion grooves along grain boundaries can be seen in the specimen serviced for 36891h, compared to the 41000h specimen, indicating that Cr depletion zones in the 36891h specimen is more severe than that in the 41000h specimen. Cu-rich phase will precipitate in Super304H steel when servicing at 600 °C. The Cu-rich phase, spherical in shape with a size of tens of nanometers, distributes evenly and dispersively in matrix. Studies show that Cu-rich phase has excellent microstructure stability and its particle size will not change obviously during long time service [12–13]. Stable Cu-rich phase in Super304H does not lead to its magnetic transition during high temperature service. Regular precipitates do not lead to magnetic transition in Super304H steel during service. Therefore, some ferromagnetic phase with magnetism must have formed in the 36,781 h Super304H tube.

Fig. 4. SEM morphologies of Super304H steels serviced for (a) 36891h; (b) 41000h.


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Fig. 5. DL-EPR curves of serviced Super304H steels.

Fig. 6. Corrosion morphologies of serviced Super304H steels after DL-EPR test.

4.2. Martensite in serviced Super304H steel tube Shi [11] found the existence of martensite in the serviced Super304H tube. The decrease of the content of Cr, C, Nb and other alloy elements near the grain boundaries causes the increase of Ms temperature near the grain boundary. Meanwhile, coarsening of grains will also cause the increase of Ms point [14]. Butler [14] studied the martensitic transformation in Cr depletion zones of TP304H steel, pointing out that martensitic transformation occurred around 10%Cr depletion zones at grain boundaries. The factors influencing martensitic transformation in Cr depletion zones are as follows [14]: 1) Cr concentration in Cr depletion zones. From the empirical formula of Ms point, it is known that the Ms point rises to room temperature when the Cr content in Cr depleted zones is less than 12%. Essentially, the mechanism of martensitic transformation in Cr depletion zones is similar with that of the traditional martensitic transformation; 2) orientation relationship of martensite with austenite matrix in Cr depleted zones. If grain boundaries move to a favorable orientation during aging process, nucleation of martensite will be easier; 3) stress concentration. Martensitic transformation occurs easily at the places where stress is not uniform, such as the places of slip band crossing or the grain boundary with high interfacial energy. 4.3. Effect of grain size on martensitic transformation of Super304H steel According to the Fig. 2, the grains of the 36891h specimen is significantly coarse. As a result, the channels for Cr diffusion decrease which will lead to weak compensation of Cr diffusion from matrix to grain boundaries [8]. According to the empirical formula of Ms in 18-8 type austenitic stainless steel as follows, decrease of Cr results in increase of Ms which may lead to formation of ferromagnetic martensite [14].

Ms ( °C) = 1305 − 41.7(%Cr) − 61.1(%Ni) − 33.3(%Mn) − 27.8(%Si) − 1667 (%[C + N] ) EBSD analysis was carried out in this study to determine whether or not ferromagnetic α martensite had formed near the grain boundaries of the 36891h tube with exceptionally coarse grains. As is shown in Fig. 7, microstructure of both serviced tubes is composed of face centered cubic austenite and M23C6 precipitated along grain boundaries. Different from the tube serviced for 41000h, lath-shaped α martensite (red mark) with a lath width of about 0.5 μm is observed near the grain boundary of the 36891h specimen. 6

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Fig. 7. EBSD images of grain boundary precipitates of Super304H tubes serviced for (a) 36891h; (b) 41000h. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)

The significant difference between the 36891h and 41000h Super304H tubes is grain size. The exceptional coarse grains in the tube serviced for 36891h will decrease the diffusion channels for Cr and lead to the severe Cr depletion and increase of Ms near grain boundaries. If the coarse grains have caused the Ms of the austenite matrix near grain boundaries to be higher than room temperature, martensitic transformation may be induced and ferromagnetic “α” phase may form during shutdown maintenance (cooling process). 5. Conclusion Conventional precipitation phases in Super304H steel did not cause obvious change in the magnetism of the serviced Super304H steel. The exceptional coarse grains in 36891h serviced tube reduced Cr diffusion channels so that Cr content in the Cr-depleted zones around carbide precipitates at grain boundaries could not be compensated by the diffusion of Cr from austenitic matrix and maintained at a lower level, which led to Ms rise in the Cr depleted zones. Martensitic transformation might then be induced during cooling process and ferromagnetic martensitic phase would form. Therefore, the main reason for the significant magnetism increase of 36891h serviced tube is attributed to its exceptional coarse grains. Acknowledgements The authors would acknowledge the financial support from the National Natural Science Foundation of China (51471072) and China Southern Power Grid Corp funded science and technology project (K-GD2014-162). References [1] I. Sen, E. Amankwah, N.S. Kumar, et al., Microstructure and mechanical properties of annealed SUS 304H austenitic stainless steel with copper[J], Mater. Sci. Eng. A 528 (13) (2011) 4491–4499. [2] Y. Sawaragi, K. Ogawa, S. Kato, et al., Development of the economical 18-8 stainless steel (SUPER304H) having high elevated temperature strength for fossil fired boilers[J], Sumitomo Search (Japan) 48 (1992) 50–58. [3] Y. Sawaragi, N. Otsuka, H. Senba, et al., Properties of a new 18-8 austenitic steel tube (SUPER304H) for fossil fired boilers after service exposure with high elevated temperature strength[J], Sumitomo Search 56 (1994) 34–43. [4] A. Ohtomo, Magnetic Measurement of Internal Scale in Austenitic Stainless Steel Tubes[R], Tokyo Ishikawajima-Harima Heavy Industries Co, Ltd, 2000. [5] Z.Y. Yang, J. Wang, J.Y. Chen, Thermal-induced martensitic transformation in 304 austenitic stainless steel[J], Transaction of Materials and Heat Treatment 29 (1) (2008) 98–101 (in Chinese). [6] W. Liu, Z.B. Li, X. Wang, et al., Effect of strain rate on strain induced α’ martensite transformation and mechanical response of austenitic stainless steels[J], Acta Metall. Sin. 45 (3) (2009) 285–291 (in Chinese). [7] K. Mumtaz, S. Takahashi, J. Echigoya, et al., Magnetic measurements of martensitic transformation in austenitic stainless steel after room temperature rolling[J], J. Mater. Sci. 39 (1) (2003) 85–97. [8] V. Mertinger, E. Nagy, F. Tranta, et al., Strain-induced martensitic transformation in textured austenitic stainless steels[J], Mater. Sci. Eng. A 481 (21) (2008) 718–722. [9] D. Jorge-Badiola, A. Iza-Mendia, I. Gutiérrez, Study by EBSD of the development of the substructure in a hot deformed 304 stainless steel[J], Mater. Sci. Eng. A


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