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308L weld and oxide films formed in deaerated high-temperature water
Journal of Nuclear Materials 498 (2018) 227e240
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Characterization of microstructure of A508III/309L/308L weld and oxide films formed in deaerated high-temperature water Qi Xiong a, b, c, Hongjuan Li a, b, c, Zhanpeng Lu a, b, c, *, Junjie Chen a, b, Qian Xiao a, b, Jiarong Ma a, b, Xiangkun Ru a, b a b c
Institute of Materials, School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road, P.O.Box 269, Shanghai, 200072, China State Key Laboratory of Advanced Special Steels, Shanghai University, 149 Yanchang Road, Shanghai, 200072, China Shanghai Key Laboratory of Advanced Ferrometallurgy, Shanghai University, 149 Yanchang Road, Shanghai, 200072, China
h i g h l i g h t s Surface films at LAS-309L/308L weld in deaerated high-temperature water at 290 C. Less surface oxides on decarburization zone than on other zones in the LAS HAZ. More corrosion pits on 309L than on 308L due to the inclusion dissolution. Film on 309L with ~9% ferrite was thicker than that on 308L with ~12% ferrite. Cr content in the inner oxide layer on 308L was higher than that on 309L.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 August 2017 Received in revised form 5 October 2017 Accepted 11 October 2017 Available online 21 October 2017
The microstructure of A508III/309L/308L weld clad and the properties of the oxide films formed in simulated pressurized water reactor primary water at 290 C were characterized. The A508III heataffected zone (HAZ) consisted primarily of a decarburization zone with ferrite near the fusion line and a following pearlite structure with fine grains. A high hardness region in the HAZ could be the result of Cenrichment. M23C6 and M7C3 precipitates were observed in element transition zone. 308L stainless steel (SS) containing ~ 12% ferrites exhibited both ferritic-austenitic solidification mode (FA mode, d/g) and austenitic-ferritic solidification mode (AF mode, g/d), whereas 309L SS containing ~ 9% ferrites exhibited only FA mode. The A508III surface oxide film was mainly Fe3O4 in deaerated high-temperature water. The coarse grain zone covered with few oxide particles was different from other types of film on the other region of HAZ and the bulk zone. More pitting appears on 309L SS after immersion in deaerated high-temperature water due to the dissolution of inclusions. SS surface oxide films consisted primarily of spinels. The oxide film on SS was divided into two layers. Ni was concentrated mainly at the oxide/ substrate interface. The oxide film formed on 309L was thicker than that on the 308L. The ferrite in the stainless steel could improve the oxidation resistance. © 2017 Elsevier B.V. All rights reserved.
Keywords: Microstructure Ferrite Low alloy steel Stainless steel Oxide film High temperature water
1. Introduction In pressurized water reactor (PWR) nuclear power plants, low alloy steels (LASs) are widely used in reactor pressure vessels, pressurizer vessels, and other vessel components because of their
* Corresponding author. Institute of Materials, School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road, P.O.Box 269, Shanghai, 200072, China. E-mail address: [email protected] (Z. Lu). https://doi.org/10.1016/j.jnucmat.2017.10.030 0022-3115/© 2017 Elsevier B.V. All rights reserved.
high strength, good ductility and low price. Stainless steels are cladded on the inner wall of the LAS pipes or vessels due to their good corrosion resistance [1e4] to mitigate the corrosion of LAS by high-temperature water coolants. Since the stainless steel cladding is in direct contact with high temperature water in the PWR primary loop, LAS would be subject to corrosion by high temperature water once the cladding layer fails. Stress corrosion cracking is one of the potential degradation modes for stainless steel cladding, which is strongly affected by the alloy microstructure and the properties of oxide films. The corrosion resistance and integrity of the stainless steel cladding are critical for the safety and reliability
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of LAS vessels and pipes. The reheat cracking in as-welded austenitic stainless steel components operating at high temperature in nuclear power plant have been reported [5]. Failures of dissimilar metal weld safe end nozzles due to stress corrosion cracking in PWR primary water have also been reported [6,7]. The material properties in the weld heat-affected zones and the weld metals in stainless steel welds and Ni-base alloy welds have been characterized [8e12] and were correlated to stress corrosion cracking behavior in high-temperature water environments. The microstructural and chemical properties in dissimilar welds would have significant effects on corrosion and stress corrosion cracking behavior in high-temperature water environments. One part of the dissimilar metal weld safe end is the stainless steel cladding on the inner wall of low alloy steel pipes. Ming et al. [13] studied the microstructural property of SA508-309L/308L-316L dissimilar metal welded safe-end joints, and found that the HAZ with a complex structure gradient appeared in LAS near the fusion boundary. The transition zone has a complex structure and carbon migration from the LAS to SS results in a hardness decrease in LAS and an increase in the carbon-rich zone [14]. Li et al. [15,16] studied the structure of low alloy steel to stainless steel transition weld and SCC behavior in simulated PWR primary water at 290 C and found that the transition zone showed a higher susceptibility to SCC. Intergranular cracking mainly occurred in the austenitic layer, whereas transgranular cracking mainly occurred in low alloy steel. The properties of oxide film have proven to play an important role in SCC initiation and the growth of materials in hightemperature water environments [17e19]. Characterization of the oxide films formed on different parts of the LAS-SS cladding in simulated primary water is expected to provide mechanistic information on SCC susceptibility. Cunha Belo et al. [20] analyzed the oxide films that formed on 316L stainless steel in high-temperature water at 350 C and found that the outer oxide film was mainly Ni0.75Fe2.25O4 inverse spinel, whereas the inner oxide film was mixed chromium oxides. Chen et al. [8,9] found that the doublelayered structure was composed of mixed iron-nickel and chromium oxides. With a decreasing the oxygen content in hightemperature water, the chromium content increased in the oxide film on stainless steel and hematite tended to diminish, whereas spinel appeared [21]. Das et al. [22] studied the oxidation behavior of Fe-Cr-Ni ternary alloys by means of quantum chemical molecular dynamics at 288 C and found that iron and chromium migrated faster in the surface layer than nickel did. Higher chromium content led to the formation of more ferrite and of compact oxide film [23]. To reveal the properties of oxide films on LAS-SS cladding, the properties of oxide films formed on various locations of the LAS-SS weld cladding block in simulated PWR primary water are characterized and correlated to the material properties at these locations. Metallographic examination, scanning electron microscopy (SEM) measurement, Raman spectroscopy, focused ion beam (FIB) sampling and transmission electron microscopy (TEM) are used in the analysis. 2. Experimental 2.1. Material and specimen preparation This work was conducted on LAS to SS cladding (A508III/309L/ 308L) that was obtained from a PWR pressure vessel nozzle and pipe weld joint mock up. Low alloy steel base metal was made from MnNi-Mo forgings, and the main chemical compositions of each part are listed in Table 1. Considering the elements (Cr, Ni) dilution of dissimilar metal welding, the first 309L SS cladding layer of about 2.5mm-thick on LAS was used as the isolation layer and proper heat treatment was conducted; then, a layer of 308L SS was used as the
corrosion resistant layer, and proper post weld heat treatment was carried out. The heat treatment parameters are shown in Table 2. The specimens with a size of 20 mm 8 mm 3 mm were cut from a plate sample. Fig. 1 displays the macroscopic sketch map of the plate, where the red box shown in Fig. 1(b) indicates the sampling location where the A508III/309L/308L welds were acquired as shown in Fig. 1(c). To reveal the microstructure, the samples were prepared by grinding and polishing. Before optical microscope observations, various parts of dissimilar metal joints were etched separately; the low alloy steel side was etched by 4% nital solution for approximately 10 s, and the stainless steel was electro-etched in 40% NaOH solution for 30 s with a DC of 5e6 V. The microstructural features were then examined using a VHX100 digital microscope and a field-emission scanning electron microscope equipped with an electron backscattered diffraction system operating at 20 kV. The fusion zone was studied using a JEM 2010F transmission electron microscope with selected area electron diffraction analyses (SAED) operating at 200 kV. To reveal the chemical composition, energy dispersive spectroscopy (EDS) equipped on an Apollo 300 SEM was used at 15 kV. 2.2. Immersion tests in high temperature water environment Immersion tests in simulated pressurized primary water environments at 290 C were conducted to study the characteristics of oxide film on the weld. The immersion samples were punched with a diameter F2 circular hole in the low alloy steel side to facilitate hanging in the autoclave. Then, the specimens were polished with metallographic sandpaper to 1500 grit before being cleaned with deionized water and alcohol. Immersion tests were carried out in deaerated PWR water. The solution for immersion tests was made by adding 2 ppm lithium (with LiOH$H2O) and 1200 ppm boron (with H3BO3) to deionized water. Deaerated water was achieved by purging N2 gas for 2.5 h and slowly heating the water to 290 C at ~7.2 MPa. For this work, the oxygen content and hydrogen content in the deaerated PWR water was expected to be below 10 ppb through bubbling with nitrogen gas into the autoclave for a period of time, which was significantly lower than that in oxygencontaining water. Deaerated high temperature water was used to simulate the local or abnormal conditions in the plants where dissolved hydrogen concentration can be low. Immersion tests were carried out for 168 h. Samples were cleaned with deionized water, rinsed in acetone and dried with cold air before using. 2.3. Characterization of oxide films formed in high temperature water The surface oxide film morphology of specimens after immersion tests was observed by Apollo 300 thermal field emission SEM. The structure of surface oxide was detected using an INVIA Raman spectrometer with a laser wavelength of 514.5 nm. Cross sections of the oxide films were used for further study with TEM observation. To prepare the TEM specimens, FEI Helios Nanolab 600i dual-beam focused ion beams with Ga ion sputtering were used to obtain the oxide film cross section. SAED was conducted using a JEOL JEM2100F TEM equipped with an EDS system operating at 200 kV. The chemical composition of oxide films formed on specimens was analyzed by EDS. 3. Results 3.1. Microstructure at various locations in the weld The microhardness profiles of the LAS-SS weld are shown in Fig. 2. The highest hardness is detected near the fusion line. The
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Table 1 Chemical compositions of the base metal and the clad metal (wt. %). Alloy
Cr
Ni
Mn
Si
C
S
Mo
Cu
Nb
Fe
A508Ⅲ 309L 308L
0.16 23.94 20.38
0.82 13.48 10.32
1.40 1.67 1.76
0.23 0.40 0.37
0.192 0.016 0.020
0.0007 <0.001 0.002
0.491 0.031 0.022
0.039 0.044 0.029
<0.001 0.001 0.001
Balance Balance Balance
Table 2 The conditions of post weld heat treatment. Temperature Time
595e620 C ~40h
hardness curve of the low alloy steel is unstable in repeated measurements, and the hardness of the SS shows a non-obvious change with an increasing distance from the fusion line.
3.1.1. Low alloy steel side Fig. 3(a) shows the optical microstructure of low alloy steel A508III. The bulk metal consists mainly of upper bainite as shown in Fig. 3(b). The HAZ close to the fusion line in A508III with a complex structure is formed and confirmed to have a width of approximately 5 mm width as shown in Fig. 3(a). The HAZ is divided into intercritically reheated coarse-grained zone (ICCGZ), fine-grained zone (FGZ) and coarse-grained zone (CGZ), as displayed in Fig. 3(c)-(e). The formation of CGZ is due to the high temperature, and grains develop in the surfacing process. The FGZ
is formed by recrystallization. The decarburization zone, which almost overlaps the CGZ by a width of approximately 500 mm, consists primarily of ferrite. The FGZ is comprised of pearlite that exists in the range of 0.5 mme5 mm from the fusion line. It can be observed that clear banded black areas appear in the HAZ but not in the bulk base metal. Fig. 4 displays the EBSD results of A508III, the grain orientation coloring maps of ICCGZ, FGZ, CGZ and the bulk base metal are shown in Fig. 4(a), with X indicating the distance from the fusion line, where the grains in each zone exhibit random orientation. Fig. 4(b) shows the corresponding grain boundary of each zone. The boundaries are mainly random grain boundaries (RGB), and some CSL grain boundaries are mainly S3.
3.1.2. Stainless steel cladding side Fig. 5 shows a metallograph of 309L SS and 308L SS. Inclusions with a diameter of approximately 18 mm are found in 309L SS cladding by metallographic observations, which are classified as SiO2 type or MnS type by SEM-EDS analysis (shown in Table 3). Obvious inclusions are not observed in 308L SS obviously.
Fig. 1. Sampling process (a) pressure vessel nozzle and pipe joint (b) sampling location (c) experimental sample.
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Fig. 2. Microhardness profiles of the LAS-SS weld from multiple measurements.
The microstructure of 309L SS near the fusion line is shown in Fig. 6. The structure of 309L SS near the fusion line is an austenite (A) zone with a 180e200 mm width. The other region of 309L/308L SS consist of a mixed austenite and ferrite (A þ F) structure, as shown in Fig. 6. Primary and secondary boundaries appear near the fusion line, where the second boundary is parallel to the fusion line and the primary boundary exists between the fusion line and second boundary perpendicular to the fusion line. Fig. 7(a) shows the morphology picture of SS cladding and the interface of 309L SS and 308L SS is clear due to the different content and form of ferrites in 309L SS and 308L SS. Using the metallographic method, the statistical results show that the ferrite contents in 309L SS and 308L SS are approximately 9% and 12% respectively. Two different forms of ferrite existed in 308 SS, as shown in Fig. 7(c) and (d). The SEM of ferrite in 309L SS and 308L SS are displayed in Fig. 7 (e) and (f), respectively, precipitates with a
diameter of approximately 500 nm that exhibit Fe and Cr-rich by EDS detection exist in ferrite phase of 309L SS, whereas obvious precipitates are not observed in the ferrite of 308L SS. Fig. 8 shows the EBSD results of 309L SS and 308L SS. The grain size within a distance of 500 mm from the fusion line is much smaller than the bulk 309L SS grain size as shown in Fig. 8(a). The bulk 309 SS and 308L SS consist primarily of columnar crystal that is perpendicular to the fusion line along the direction of heat dissipation, and the grains show [001] preferred orientation, as displayed in Fig. 8(c). The CSL boundary in SS occurs mainly near the fusion line, as shown in Fig. 8(b). 3.1.3. Interface of LAS-SS The element profiles near the A508III-309L SS fusion line are shown in Fig. 9. The concentrations of the main elements Fe, Cr, and Ni change rapidly near the fusion line, Cr and Ni migrate from 309L
Fig. 3. Optical micrographs of A508III, (a) the low magnification metallographic panorama, (b)e(e) the enlarged organization image of the corresponding position in (a) respectively.
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Fig. 4. EBSD results of A508III: (a) grain orientation coloring maps and (b) grain boundary character maps. X indicates the distance from the fusion line.
Fig. 5. Metallographs of 309L SS and 308L SS after polishing.
SS to A508III and Fe diffuses in the opposite direction, as shown in Fig. 9(a). Carbon migration has been reported by many researchers [20,21]. Carbons migrates from A508III to the 309L SS side, forming a carbon-depleted region, as shown in Fig. 3(a). The migration of carbon is due to the different carbon contents of dissimilar metal and chromium is an element that possesses a strong affinity for carbon. Fig. 9(b) shows the details of element migration, the change of elements occurred in the region (called element transition zone) of about 100 mm width. Fig. 10 shows the TEM and SAED results of the region near the fusion line. The sampling position is shown in Fig. 10(a), which was obtained by FIB. The SAED results indicate the 309L SS substrate near the fusion line has an austenite structure (FCC), as shown in Fig. 10(d), and the element transition zone matrix shows a structure with BCC in Fig. 10(e). Fig. 10(c) displays a TEM image of the element transition zone and granular or rod-shaped precipitates
are found in this region. Fig. 10(f) shows the SAED result, which indicates that those precipitated pellets are mainly metallic carbide. Both M23C6 and M7C3 (M is mainly Fe and Cr) that show an orientation relationship of axis [01-1]//[331] is detected.
Table 3 SEM-EDS results of the chemical compositions (wt. %) of inclusions and pits on 309L SS after immersion in deaerated high temperature water.
Inclusions Pits
Fe
Cr
Ni
Mn
Si
O
S
9.41 22.22
6.50 22.55
0.81 1.92
32.42 7.56
17.86 0.63
27.09 20.48
1.10 0.19
Fig. 6. Optical micrograph of stainless steel near the fusion line.
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Fig. 7. Optical micrographs of (a)309L SS and 308L SS, (b) ferrites in 309L SS, (c) (d) ferrites in 308L SS, (e)(f) SEM maps of ferrites in 309L SS and 308L SS respectively.
3.2. Oxide films formed on the specimen surface after immersion in deaerated high temperature water at 290 C After the immersion test for 168 h, oxide films formed on the weld specimen are characterized. Fig. 11 shows the low
magnification of film on the weld surface. The interface of oxide film formed at various parts is obvious. The region of A508III close to the fusion line shows a bright zone, similar to the decarburization zone. Corrosion pits are detected on 309L SS, as displayed in Fig. 11(c), but obvious pits are not found on the 308L SS surface.
Fig. 8. EBSD results of 309L SS and 308L SS: (a) grain orientation coloring maps (b) grains boundary character maps and (c) inverse pole maps. X indicates the distance from the fusion line.
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Fig. 9. SEM-EDS results (a) Distribution of some elements near the A508III-309L SS fusion line and (b) details.
Fig. 12 shows the SEM images of pits on 309L SS, and the average size of pits is approximately 20 mm in diameter. Fig. 13 shows SEM images of the oxide film formed on the A508III in the weld. Fig. 14 show SEM images of the oxide film formed on stainless steel. Fig. 15 shows the Raman spectroscopy results obtained for the oxide films formed on LAS and SS after immersion in high temperature water, where N1-N6 in the figures indicates the locations far from the fusion line, and the corresponding position (except N1) is shown in Fig. 11. N1, N2 and N3 are located in the bulk of A508III, FGZ and CGZ, respectively, whereas N4 is located in austenite zone, and N5 and N6 are located in 309L and 308L, respectively. The TEM results of the oxide films formed on 309L SS and 308L SS surfaces after immersion in deaerated hightemperature water are shown in Figs. 16e21.
3.2.1. Oxide films formed on A508 III The surface of CGZ is covered with almost no oxide particles, but those on FGZ, ICCGZ and the bulk zone are no different according to the SEM observation, as shown in Fig. 13(c)-(d). The Raman spectra
indicate that the types of the oxide films formed on various regions of A508 III are similar and that the oxide film consists primarily of Fe3O4, as shown in Fig. 15(a). 3.2.2. Oxide films formed on 309L/308L Relatively small and sparse particles are distributed in the fusion zone, as shown in Fig. 14(a). The surfaces of 309L and 308L SS are covered mainly with faceted oxide particles, as shown in Fig. 14(b) and (c). The faceted particles are composed primarily of Fe-rich oxides with low Cr and Ni contents, and the EDS results are shown in Table 4. Considering the low concentrations of oxygen in the oxide film in Table 4, it is thought that the compositions obtained by SEM-EDS were probably influenced by base metal underneath the oxide film. The Raman peaks are characteristics of spinel with AB2O4 on both the 309L SS and 308L SS surfaces, according to the characteristic peaks [24], and the spinel oxides should be (Fe,Ni) (Fe,Cr)2O4 according to EDS results. The oxide particles on 309L SS are denser in distribution than are those on the 308L SS surface.
Fig. 10. (a) Sampling location by FIB for TEM observation, (b) 309L SS substrate near the fusion line, (c) Element transition zone matrix, (d) SAED pattern of stainless steel substrate, (e) SAED pattern of element transition zone e matrix (f) SAED pattern of M23C6 (zone axis [01-1]) and M7C3 (zone axis [331]).
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Fig. 11. Low magnification photos of immersion sample in deaerated water (a) low image panorama (b)interface (c)309L (d)308L.
4. Discussion 4.1. Microstructure
Fig. 12. SEM images of pits on 309L SS after 168 h immersion in deaerated hightemperature water at 290 C.
Fig. 16 show the cross section images of the oxide films on 309L and 308L, showing a thicker inner oxide film on 309L than on 308L. The oxide films show a duplex structure on both 309L SS and 308L SS. The outer layer is composed of large and dispersive oxide particles that are determined to be spinel-type structures by the diffraction pattern (as shown in Figs. 17(I) and 19(I)). The inner layer of the oxide film on 309L SS is 140 ± 20 nm thick (region II in Fig. 17), whereas the inner layer of the oxide film on 308L SS is 60 ± 30 nm thick (region II in Fig. 19). The SAED pattern of region II from the inner layer showed diffraction rings which indicate the presence of nanocrystalline spinel oxides in the inner layer of 309L SS and 308L SS, and the dark area (region III in Figs. 17 and 19) is a matrix with FCC structure, which indicates that the matrix of 309L SS and 308L SS is shows an austenite structure. Figs. 18 and 20 show EDS mappings of the oxide layers formed on 309L SS and 308L SS respectively. The EDS results show that the outer oxide layer on both 309L SS and 308L SS is Fe-rich and that the inner layer is Crrich; further, Ni is enriched in the interface between the oxide and the matrix. The character of the element distribution is revealed in the line scan profiles, as shown in Fig. 21, which agree with the result of EDS mappings of the oxide films.
According to metallographic observations after microhardness tests (shown in Fig. 2), the unstable hardness in the HAZ is due to the black stripe, which shows a higher hardness, and the black stripe should be carbon condensed. The highest hardness near the fusion line may relate to a martensite structure. In fact, increasing carbon and chromium contents favors the formation of martensite [25], and many researchers have found that martensite exists near the fusion zone in dissimilar joints [13,14]. Based on the Schaeffler diagram, martensite will appear at a certain dilution, but in this work martensite structure is not found by TEM, possibly due to sampling position or discontinuous distribution of martensite. A similar primary boundary and secondary boundary relation was also found in other studies [26]. It is worth noting that the distribution of the secondary boundary is discontinuous. Nelson et al. [27] suggested that the primary boundary is a result of epitaxial growth based on the existing substrate during solidification. The formation of the secondary boundary depends mainly on the change in the solidification mode during the solidification process [28]. Molten metal solidification near the fusion line is ferrite (BCC) and then austenite (FCC) solidification in 309L SS, with a change in solidification. Nelson et al. considered the formation of the secondary boundary is to be the result of grain boundary migration during austenite solidification. Yoo et al. [29] studied the effects of thermal aging on the microstructure of Type-II boundaries in Ni-based and low alloy steel dissimilar welds and found that the Type-II boundaries shifted away from the fusion line with an increase in the thermal aging time at 450 C. Therefore, the secondary boundary should be formed by the migration of the austenite boundary. Previous works have found that the secondary boundary, which is parallel to fusion boundary in the filler metal in dilution zone is a potential crack path for SCC [27], which will speed up the failure of the weld. The bulk SS shows a mixed austenite and ferrite structure, and the ferrites in 309L SS and 308L SS show differences in content and shape. From the Schaeffler diagram [30], the ratio of Creq/Nieq in 309L/308L SS falls into the gþd region. Cr is a ferrite-stable element, and a high Cr content is favorable for the formation of ferrite, whereas a higher content of Ni, which is an austenite-stable element in 309L SS, has a negative effect on the stability of ferrite. As shown in Table 5 from SEM-EDS analyses, the ferrites
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Fig. 13. SEM images of oxide film on LAS after 168 h immersion in deaerated high-temperature water at 290 C, (a) film on CGZ (b) film on FGZ (c) film on ICCGZ (d) film on bulk base metal of A508III.
Fig. 14. SEM images of oxide film after 168 h immersion in deaerated high temperature water at 290 C, (a)interface (b)309L (c)308L.
distributed in SS reveal a higher ratio of Cr/Ni than austenite does. Heat input is also crucial to the ferrite content [31], increasing the welding temperature and heat treatment time decreases the content of ferrite. Heat treatment after cladding 309L SS and cladding the second layer of 308L SS are also factors that cause 309L SS to exhibit a higher ferrite content than 308L SS. Relevant research has found that the distribution of ferrite in austenite stainless was related to the ratio of Creq/Nieq and welding speeds which determine the solidification mode of Fe-Cr-Ni alloy [31e33]. A primary ferrite following d transform to g (FA mode) or a primary austenite following g transform to d upon cooling (AF mode) may occur on austenitic stainless steel welds. Chemical composition and welding process conditions determine solidification mode. The following
four solidification modes have been reported by previous researchers [34]: F mode: Creq/Nieq (in wt.%)< 1.25, Lþ g/g
(1)
AF mode: 1.25< Creq/Nieq (in wt.%) < 1.48, Lþ g/Lþ gþd/gþd
(2)
FA mode: 1.48< Creq/Nieq (in wt.%) < 1.95, Lþd/Lþdþg/dþg
(3)
A mode: 1.95< Creq/Nieq (in wt.%), Lþd/d
(4)
where
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Fig. 15. Raman spectra of the oxide films on weld after 168 h immersion in deaerated high-temperature water at 290 C, (a) A508 III (b) 309L/308L SS.
Fig. 16. TEM images of oxide films on (a) 309L SS (b) 308L SS after 168 h immersion in deaerated high-temperature water at 290 C.
Fig. 17. TEM image for area 1 in Fig. 16a) and correspondent electron diffraction patterns of oxide film on 309L SS after exposure in deaerated water. The white dotted line indicates the interface of the outer-layer and inner-layer oxide film.
Creq ¼ Cr þ Mo þ 1.5 Si þ 0.5 Nb
(5)
Nieq ¼ Ni þ 30 C þ 0.5 Mn
(6)
For the present work, the Creq/Nieq of 309L SS is 1.66, and the Creq/Nieq of 308L SS is 1.77. It is suggested that the solidification of
309L SS and 308L SS falls into the FA mode. However, both AF mode and FA mode are observed on 308L SS as shown in Fig. 7(a). As observed in other studies [35e37], FA mode presents a greater amount of ferrite with lathy, whereas AF mode presents skeletal ferrite, similar to the results shown in Fig. 7(c) and (d). Suutala [33] reported that the welding speed and cooling rate have an effect on
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Fig. 18. STEM-HAADF image of oxide film on 309L SS and the corresponding EDS mappings for Fe, Cr, Ni and O after immersion in deaerated water.
Fig. 19. TEM image for area 2 in Fig. 16b) and correspondent electron diffraction patterns of the oxide film on 308L SS after exposure in deaerated water. The white dotted line indicates the interface of the outer-layer and inner-layer oxide film.
Fig. 20. STEM-HAADF image of oxide film on 309L SS and the corresponding EDS mappings for Fe, Cr, Ni and O after immersion in deaerated water.
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Fig. 21. EDS line scan profiles of the oxide film in deaerated water, (a) on 309L SS along the white dotted line a as shown in Fig. 15, (b) on 308L SS along the white dotted line b shown in Fig. 17.
Table 4 SEM-EDS results of the chemical compositions (wt. %) of the oxide particles (shown in Fig. 14) after immersion in deaerated high temperature water.
Particles on 309L SS Particles on 308L SS
Fe
Cr
Ni
O
67.78 60.74
4.74 8.09
2.47 4.39
22.70 22.22
Table 5 SEM-EDS results of austenite and ferrite in 309L SS cladding alloy. Phases
Fe
Cr
Ni
Cr/Ni
Austenite Ferrite
69.20 73.38
17.02 20.61
10.48 3.18
1.62 6.48
the solidification mode. The steady-state dendritic growth indicates the effect of the solidification conditions and a higher steady-state dendritic growth rate favors austenite as the primary phase. A high welding speed and cooling rate are preferred for AF mode. In the weld pool, the temperature gradient and element segregation during solidification are the main factors causing AF and FA mode to occur in the 308L SS surface layer.
4.2. Effects of LAS microstructure on the oxide film formed in deaerated high-temperature water Fe3O4 particles are detected on the LAS surface film in deaerated water, as shown in Fig. 15(a). From the potential-pH diagram [38,39], Fe3O4 is more stable at low potential whereas Fe2O3 is more stable at high potential, which suggests that Fe3O4 is easier to form in deaerated high-temperature water. Meisel [40] predicted the iron oxide or iron hydroxide type by calculating the effective Gibbs free energy Geff, where the generated Geff of the nonstoichiometric compound Fe3O4 is the minimum value. Fe3O4 should be the product of LAS corrosion in deaerated high-temperature water. The Raman spectra show that there is no significant difference in oxide species on each region of the LAS surface, whereas the surface oxide distribution on CGZ is quite different from that on FGZ, ICCGZ and the bulk metal, as displayed in Fig. 13. The FGZ consists mainly of ferrite coinciding and has a lower carbon content than do other regions of LAS due to decarburization. Liu et al. [41] studied the influence of carbon content on corrosion resistance of low alloy steel and found that the corrosion potential of corrosion layer becomes more positive, the corrosion current density
decreases and the oxide film has more protection performance with a decreasing carbon content of LAS. Clover et al. [42] considered that the main reason why the pearlite phase is worse than ferrite in corrosion resistance is that pearlite is a eutectoid body of ferrite and cementite; further, as the potential of the two-phase is different, there is severe micro galvanic corrosion in pearlite phase and cementite because the cathode will accelerate the corrosion of ferrite in pearlite. Therefore decarburization may be the main reason for the CGZ showing almost no particles on the film compared to other regions. 4.3. Effects of stainless steel microstructure on the oxide film formed in deaerated high-temperature water The oxygen concentration has a great effect on the electrochemical corrosion potential of metal in high-temperature water [43]. Potential-pH diagrams predicted that the spinel-type oxide is more stable on Fe-Cr-Ni alloy in low-oxygen concentration water. Stellwag [44] suggested that the outer oxide film is formed by the dissolution of metal ions in deposition and that the rate of Fe atoms migration is faster than those of Ni and Cr in oxide. Thus, the outer oxide particles are Fe-enriched, whereas Cr and Ni generally remain in the inner layer. The oxide films formed on 309L SS and 308L SS were observed by SEM and TEM. Combined with the EDS results, the surface oxide particles are Fe-rich spinel, consistent with the Raman spectroscopy results. Kuang et al. [45] studied the effect of the dissolution of Ni on the oxidation of 304 SS at 290 C and found that spinel particles started to appear when low Ni2þ concentration transferred to high in solution, which suggests that Ni is favorable for spinel nucleation. 309L SS matrix has a higher Ni content. Therefore, the spinel-type particles distributed on 309L SS are denser than those distributed on 308L SS surfaces. The oxidation rate of Ni is lower than those of Fe and Cr in hightemperature water. Because of its low oxygen affinity, Ni is the least likely to be oxidized compared with Fe and Cr [46,47]. Researchers have found Fe and Cr atoms tend to deviate from the surface more easily than do Ni atoms and that Ni enrichment appeared near the inner oxide film in studies on stress corrosion behavior of Fe-Cr-Ni alloy that used the molecular dynamics model [22]. Therefore, the migration rate of Ni into the oxide film is very low, resulting in an enrichment of Ni in the interface of the inner layer and matrix, which can be clearly observed in the EDS mapping. The TEM results show that the oxide inner layer on 309L SS is thicker than that on 308L SS. EDS mapping and line scan profiles
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both reveal that the inner oxide layer is consisted of Cr-rich oxides. The enrichment of Cr in the inner layer contributed to the corrosion resistance. The oxidation rate constant of Ni-based alloys in simulated PWR primary water was calculated [17,18,48], and the oxidation rate constant of alloys 600 and 182 with 15% Cr content was higher than that of alloy 690 with 30% Cr content. Das et al. [22] found that atoms diffused slower in Cr-rich oxide than in Fe-rich and Ni-rich oxide and that a Cr-rich oxide layer is counterproductive to the growth of oxide film. The inner layer of the oxide film formed on 316L SS with a higher Cr content was thinner than that on 9Cr-26Ni-1.5Mo-5Si alloy as reported in Ref. [45]. The oxidation rate of stainless steel in high-temperature water is mainly determined by the Cr-rich inner oxide layer. Data from multiple measurements indicate that the content of Cr of the inner oxide layer on 309L SS is 24.4% (wt%) whereas the content of Cr in the inner oxide layer on 308L SS is 33.47% (wt%). The growth of oxide film depends on the metal atoms migrating from matrix to oxide film and oxygen diffusing into oxide film. The oxide film with a higher Cr content can mitigate more effectively the transport of metal atoms and oxygen atoms, which is supposed to result in a lower oxide film. However, experimental data show that Cr content of 309L SS substrate is higher than that of 308L SS substrate. The higher Cr content in the inner oxide layer of 309L SS can be related to the effect of ferrite on oxidation. Cao et al. [23] found that higher Cr content led to a higher ferrite content and reduced the amount of oxidized metal. Post weld heat treatment (PWHT) had no significant effect on the formation of oxide film. Their work did not emphasize the effect of ferrite on the oxide formation. Previous works have reported that the interaction between the austenite and the ferrite would affect the corrosion potential of dual-phase steel [49]. The galvanic interaction influence of a-phase in stainless steel has been investigated [50,51]. a-phase could move the open circuit potential of g-phase toward the noble direction, and the passive film impedance of single g-phase increased after the contact of the two phases. The oxide film on the duplex stainless steel showed better corrosion resistance. Lu et al. [52] reported that the crack growth rates in the weld metal with ~6.5 wt% ferrite content were higher than those in the weld metal with ~8.5 wt% ferrite content in dissolved oxygen water at 288 C. Those studies indicated that the ferrite content of SS affects the property of oxide film formed on SS. Compared with 309L SS, 308L SS has a higher ferrite content but a lower Cr content. Ferrites in stainless steel may improve the Cr content of oxide film, which indicates a lower growth rate of oxide film on 308L SS than on 309L SS. However, the detailed reasons for the effect of ferrite to Cr concentration in oxide film require further study. Many corrosion pits could be found only on the 309L SS surface after immersion. Obviously inclusions are observed in 309L SS but are not observed in 308L SS before the immersion test. Through observation and statistics, the distribution and size of inclusions and pits were found to be similar. The EDS results of inclusions and pits are displayed in Table 3. The Si and Mn contents decreased sharply in pits compared with inclusions, which suggest inclusions may have dissolved in deaerated high-temperature water and formed the pits. The preferential dissolution of Si in oxide films in high-temperature water has been reported [53]. This work suggests that silicon oxide has high solubility in high-temperature water which agrees with this result. It confirms that corrosion pits on the film of 309L SS surfaces are caused by inclusion dissolution. These corrosion pits will act as channels for crack propagation and accelerate the stress corrosion cracking. Nonmetallic inclusions in steels affect corrosion resistance, such as pitting corrosion, stress corrosion cracking and hydrogen-induced cracking [54,55] resistance.
239
5. Conclusions The microstructure and properties of the oxide films formed on various locations of a A508III/309L/308L weld clad in simulated deaerated PWR high-temperature water are characterized as follows. 1. The HAZ in the LAS side consisted of CGZ, FGZ and ICCGZ. A high hardness zone in the SS zone near the fusion line was identified, and attributed to the C-enrichment. M23C6 and M7C3 precipitates in element transition zone were identified by TEM. 2. Ferrite content was approximately 9% in 309L SS and approximately 12% in 308L SS. Both FA and AF mode were found in 308L SS, but only FA mode was found in 309L SS cladding. 3. The surface oxide films on LAS HAZ and base metal in deaerated high-temperature water were mainly Fe3O4 type, where CGZ exhibited less oxide coverage than did other regions in the LAS HAZ. The oxide films on stainless steel cladding surface were mainly spinel oxides. More corrosion pits appeared on 309L SS than on 308L SS after immersion in deaerated high-temperature water due to the dissolution of inclusions. 4. The oxide film formed on SS cladding in high temperature water can be divided into two layers: the Fe-rich spinel outer oxide layer and the Cr-rich oxide inner layer. Ni-enrichment at the oxide/substrate interface is observed. The oxide film formed on 309L is thicker than that on the 308L, which is thought to be related to the effect of ferrite on oxidation. Acknowledgments This work has been supported by Natural Science Foundation of China (NSFC No. 51771107), National Key R&D Program of China (No. 2017YFB0703002), and Shanghai Municipal Commission of Economy and Informatization (No. T-221715003). The support from the Instrument Analytical and Research Center, Shanghai University is acknowledged. References [1] [2] [3] [4] [5]
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Characterization of microstructure of A508III/309L/308L weld and oxide films formed in deaerated high-temperature water Qi Xiong a, b, c, Hongjuan Li a, b, c, Zhanpeng Lu a, b, c, *, Junjie Chen a, b, Qian Xiao a, b, Jiarong Ma a, b, Xiangkun Ru a, b a b c
Institute of Materials, School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road, P.O.Box 269, Shanghai, 200072, China State Key Laboratory of Advanced Special Steels, Shanghai University, 149 Yanchang Road, Shanghai, 200072, China Shanghai Key Laboratory of Advanced Ferrometallurgy, Shanghai University, 149 Yanchang Road, Shanghai, 200072, China
h i g h l i g h t s Surface films at LAS-309L/308L weld in deaerated high-temperature water at 290 C. Less surface oxides on decarburization zone than on other zones in the LAS HAZ. More corrosion pits on 309L than on 308L due to the inclusion dissolution. Film on 309L with ~9% ferrite was thicker than that on 308L with ~12% ferrite. Cr content in the inner oxide layer on 308L was higher than that on 309L.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 August 2017 Received in revised form 5 October 2017 Accepted 11 October 2017 Available online 21 October 2017
The microstructure of A508III/309L/308L weld clad and the properties of the oxide films formed in simulated pressurized water reactor primary water at 290 C were characterized. The A508III heataffected zone (HAZ) consisted primarily of a decarburization zone with ferrite near the fusion line and a following pearlite structure with fine grains. A high hardness region in the HAZ could be the result of Cenrichment. M23C6 and M7C3 precipitates were observed in element transition zone. 308L stainless steel (SS) containing ~ 12% ferrites exhibited both ferritic-austenitic solidification mode (FA mode, d/g) and austenitic-ferritic solidification mode (AF mode, g/d), whereas 309L SS containing ~ 9% ferrites exhibited only FA mode. The A508III surface oxide film was mainly Fe3O4 in deaerated high-temperature water. The coarse grain zone covered with few oxide particles was different from other types of film on the other region of HAZ and the bulk zone. More pitting appears on 309L SS after immersion in deaerated high-temperature water due to the dissolution of inclusions. SS surface oxide films consisted primarily of spinels. The oxide film on SS was divided into two layers. Ni was concentrated mainly at the oxide/ substrate interface. The oxide film formed on 309L was thicker than that on the 308L. The ferrite in the stainless steel could improve the oxidation resistance. © 2017 Elsevier B.V. All rights reserved.
Keywords: Microstructure Ferrite Low alloy steel Stainless steel Oxide film High temperature water
1. Introduction In pressurized water reactor (PWR) nuclear power plants, low alloy steels (LASs) are widely used in reactor pressure vessels, pressurizer vessels, and other vessel components because of their
* Corresponding author. Institute of Materials, School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road, P.O.Box 269, Shanghai, 200072, China. E-mail address: [email protected] (Z. Lu). https://doi.org/10.1016/j.jnucmat.2017.10.030 0022-3115/© 2017 Elsevier B.V. All rights reserved.
high strength, good ductility and low price. Stainless steels are cladded on the inner wall of the LAS pipes or vessels due to their good corrosion resistance [1e4] to mitigate the corrosion of LAS by high-temperature water coolants. Since the stainless steel cladding is in direct contact with high temperature water in the PWR primary loop, LAS would be subject to corrosion by high temperature water once the cladding layer fails. Stress corrosion cracking is one of the potential degradation modes for stainless steel cladding, which is strongly affected by the alloy microstructure and the properties of oxide films. The corrosion resistance and integrity of the stainless steel cladding are critical for the safety and reliability
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of LAS vessels and pipes. The reheat cracking in as-welded austenitic stainless steel components operating at high temperature in nuclear power plant have been reported [5]. Failures of dissimilar metal weld safe end nozzles due to stress corrosion cracking in PWR primary water have also been reported [6,7]. The material properties in the weld heat-affected zones and the weld metals in stainless steel welds and Ni-base alloy welds have been characterized [8e12] and were correlated to stress corrosion cracking behavior in high-temperature water environments. The microstructural and chemical properties in dissimilar welds would have significant effects on corrosion and stress corrosion cracking behavior in high-temperature water environments. One part of the dissimilar metal weld safe end is the stainless steel cladding on the inner wall of low alloy steel pipes. Ming et al. [13] studied the microstructural property of SA508-309L/308L-316L dissimilar metal welded safe-end joints, and found that the HAZ with a complex structure gradient appeared in LAS near the fusion boundary. The transition zone has a complex structure and carbon migration from the LAS to SS results in a hardness decrease in LAS and an increase in the carbon-rich zone [14]. Li et al. [15,16] studied the structure of low alloy steel to stainless steel transition weld and SCC behavior in simulated PWR primary water at 290 C and found that the transition zone showed a higher susceptibility to SCC. Intergranular cracking mainly occurred in the austenitic layer, whereas transgranular cracking mainly occurred in low alloy steel. The properties of oxide film have proven to play an important role in SCC initiation and the growth of materials in hightemperature water environments [17e19]. Characterization of the oxide films formed on different parts of the LAS-SS cladding in simulated primary water is expected to provide mechanistic information on SCC susceptibility. Cunha Belo et al. [20] analyzed the oxide films that formed on 316L stainless steel in high-temperature water at 350 C and found that the outer oxide film was mainly Ni0.75Fe2.25O4 inverse spinel, whereas the inner oxide film was mixed chromium oxides. Chen et al. [8,9] found that the doublelayered structure was composed of mixed iron-nickel and chromium oxides. With a decreasing the oxygen content in hightemperature water, the chromium content increased in the oxide film on stainless steel and hematite tended to diminish, whereas spinel appeared [21]. Das et al. [22] studied the oxidation behavior of Fe-Cr-Ni ternary alloys by means of quantum chemical molecular dynamics at 288 C and found that iron and chromium migrated faster in the surface layer than nickel did. Higher chromium content led to the formation of more ferrite and of compact oxide film [23]. To reveal the properties of oxide films on LAS-SS cladding, the properties of oxide films formed on various locations of the LAS-SS weld cladding block in simulated PWR primary water are characterized and correlated to the material properties at these locations. Metallographic examination, scanning electron microscopy (SEM) measurement, Raman spectroscopy, focused ion beam (FIB) sampling and transmission electron microscopy (TEM) are used in the analysis. 2. Experimental 2.1. Material and specimen preparation This work was conducted on LAS to SS cladding (A508III/309L/ 308L) that was obtained from a PWR pressure vessel nozzle and pipe weld joint mock up. Low alloy steel base metal was made from MnNi-Mo forgings, and the main chemical compositions of each part are listed in Table 1. Considering the elements (Cr, Ni) dilution of dissimilar metal welding, the first 309L SS cladding layer of about 2.5mm-thick on LAS was used as the isolation layer and proper heat treatment was conducted; then, a layer of 308L SS was used as the
corrosion resistant layer, and proper post weld heat treatment was carried out. The heat treatment parameters are shown in Table 2. The specimens with a size of 20 mm 8 mm 3 mm were cut from a plate sample. Fig. 1 displays the macroscopic sketch map of the plate, where the red box shown in Fig. 1(b) indicates the sampling location where the A508III/309L/308L welds were acquired as shown in Fig. 1(c). To reveal the microstructure, the samples were prepared by grinding and polishing. Before optical microscope observations, various parts of dissimilar metal joints were etched separately; the low alloy steel side was etched by 4% nital solution for approximately 10 s, and the stainless steel was electro-etched in 40% NaOH solution for 30 s with a DC of 5e6 V. The microstructural features were then examined using a VHX100 digital microscope and a field-emission scanning electron microscope equipped with an electron backscattered diffraction system operating at 20 kV. The fusion zone was studied using a JEM 2010F transmission electron microscope with selected area electron diffraction analyses (SAED) operating at 200 kV. To reveal the chemical composition, energy dispersive spectroscopy (EDS) equipped on an Apollo 300 SEM was used at 15 kV. 2.2. Immersion tests in high temperature water environment Immersion tests in simulated pressurized primary water environments at 290 C were conducted to study the characteristics of oxide film on the weld. The immersion samples were punched with a diameter F2 circular hole in the low alloy steel side to facilitate hanging in the autoclave. Then, the specimens were polished with metallographic sandpaper to 1500 grit before being cleaned with deionized water and alcohol. Immersion tests were carried out in deaerated PWR water. The solution for immersion tests was made by adding 2 ppm lithium (with LiOH$H2O) and 1200 ppm boron (with H3BO3) to deionized water. Deaerated water was achieved by purging N2 gas for 2.5 h and slowly heating the water to 290 C at ~7.2 MPa. For this work, the oxygen content and hydrogen content in the deaerated PWR water was expected to be below 10 ppb through bubbling with nitrogen gas into the autoclave for a period of time, which was significantly lower than that in oxygencontaining water. Deaerated high temperature water was used to simulate the local or abnormal conditions in the plants where dissolved hydrogen concentration can be low. Immersion tests were carried out for 168 h. Samples were cleaned with deionized water, rinsed in acetone and dried with cold air before using. 2.3. Characterization of oxide films formed in high temperature water The surface oxide film morphology of specimens after immersion tests was observed by Apollo 300 thermal field emission SEM. The structure of surface oxide was detected using an INVIA Raman spectrometer with a laser wavelength of 514.5 nm. Cross sections of the oxide films were used for further study with TEM observation. To prepare the TEM specimens, FEI Helios Nanolab 600i dual-beam focused ion beams with Ga ion sputtering were used to obtain the oxide film cross section. SAED was conducted using a JEOL JEM2100F TEM equipped with an EDS system operating at 200 kV. The chemical composition of oxide films formed on specimens was analyzed by EDS. 3. Results 3.1. Microstructure at various locations in the weld The microhardness profiles of the LAS-SS weld are shown in Fig. 2. The highest hardness is detected near the fusion line. The
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Table 1 Chemical compositions of the base metal and the clad metal (wt. %). Alloy
Cr
Ni
Mn
Si
C
S
Mo
Cu
Nb
Fe
A508Ⅲ 309L 308L
0.16 23.94 20.38
0.82 13.48 10.32
1.40 1.67 1.76
0.23 0.40 0.37
0.192 0.016 0.020
0.0007 <0.001 0.002
0.491 0.031 0.022
0.039 0.044 0.029
<0.001 0.001 0.001
Balance Balance Balance
Table 2 The conditions of post weld heat treatment. Temperature Time
595e620 C ~40h
hardness curve of the low alloy steel is unstable in repeated measurements, and the hardness of the SS shows a non-obvious change with an increasing distance from the fusion line.
3.1.1. Low alloy steel side Fig. 3(a) shows the optical microstructure of low alloy steel A508III. The bulk metal consists mainly of upper bainite as shown in Fig. 3(b). The HAZ close to the fusion line in A508III with a complex structure is formed and confirmed to have a width of approximately 5 mm width as shown in Fig. 3(a). The HAZ is divided into intercritically reheated coarse-grained zone (ICCGZ), fine-grained zone (FGZ) and coarse-grained zone (CGZ), as displayed in Fig. 3(c)-(e). The formation of CGZ is due to the high temperature, and grains develop in the surfacing process. The FGZ
is formed by recrystallization. The decarburization zone, which almost overlaps the CGZ by a width of approximately 500 mm, consists primarily of ferrite. The FGZ is comprised of pearlite that exists in the range of 0.5 mme5 mm from the fusion line. It can be observed that clear banded black areas appear in the HAZ but not in the bulk base metal. Fig. 4 displays the EBSD results of A508III, the grain orientation coloring maps of ICCGZ, FGZ, CGZ and the bulk base metal are shown in Fig. 4(a), with X indicating the distance from the fusion line, where the grains in each zone exhibit random orientation. Fig. 4(b) shows the corresponding grain boundary of each zone. The boundaries are mainly random grain boundaries (RGB), and some CSL grain boundaries are mainly S3.
3.1.2. Stainless steel cladding side Fig. 5 shows a metallograph of 309L SS and 308L SS. Inclusions with a diameter of approximately 18 mm are found in 309L SS cladding by metallographic observations, which are classified as SiO2 type or MnS type by SEM-EDS analysis (shown in Table 3). Obvious inclusions are not observed in 308L SS obviously.
Fig. 1. Sampling process (a) pressure vessel nozzle and pipe joint (b) sampling location (c) experimental sample.
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Fig. 2. Microhardness profiles of the LAS-SS weld from multiple measurements.
The microstructure of 309L SS near the fusion line is shown in Fig. 6. The structure of 309L SS near the fusion line is an austenite (A) zone with a 180e200 mm width. The other region of 309L/308L SS consist of a mixed austenite and ferrite (A þ F) structure, as shown in Fig. 6. Primary and secondary boundaries appear near the fusion line, where the second boundary is parallel to the fusion line and the primary boundary exists between the fusion line and second boundary perpendicular to the fusion line. Fig. 7(a) shows the morphology picture of SS cladding and the interface of 309L SS and 308L SS is clear due to the different content and form of ferrites in 309L SS and 308L SS. Using the metallographic method, the statistical results show that the ferrite contents in 309L SS and 308L SS are approximately 9% and 12% respectively. Two different forms of ferrite existed in 308 SS, as shown in Fig. 7(c) and (d). The SEM of ferrite in 309L SS and 308L SS are displayed in Fig. 7 (e) and (f), respectively, precipitates with a
diameter of approximately 500 nm that exhibit Fe and Cr-rich by EDS detection exist in ferrite phase of 309L SS, whereas obvious precipitates are not observed in the ferrite of 308L SS. Fig. 8 shows the EBSD results of 309L SS and 308L SS. The grain size within a distance of 500 mm from the fusion line is much smaller than the bulk 309L SS grain size as shown in Fig. 8(a). The bulk 309 SS and 308L SS consist primarily of columnar crystal that is perpendicular to the fusion line along the direction of heat dissipation, and the grains show [001] preferred orientation, as displayed in Fig. 8(c). The CSL boundary in SS occurs mainly near the fusion line, as shown in Fig. 8(b). 3.1.3. Interface of LAS-SS The element profiles near the A508III-309L SS fusion line are shown in Fig. 9. The concentrations of the main elements Fe, Cr, and Ni change rapidly near the fusion line, Cr and Ni migrate from 309L
Fig. 3. Optical micrographs of A508III, (a) the low magnification metallographic panorama, (b)e(e) the enlarged organization image of the corresponding position in (a) respectively.
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Fig. 4. EBSD results of A508III: (a) grain orientation coloring maps and (b) grain boundary character maps. X indicates the distance from the fusion line.
Fig. 5. Metallographs of 309L SS and 308L SS after polishing.
SS to A508III and Fe diffuses in the opposite direction, as shown in Fig. 9(a). Carbon migration has been reported by many researchers [20,21]. Carbons migrates from A508III to the 309L SS side, forming a carbon-depleted region, as shown in Fig. 3(a). The migration of carbon is due to the different carbon contents of dissimilar metal and chromium is an element that possesses a strong affinity for carbon. Fig. 9(b) shows the details of element migration, the change of elements occurred in the region (called element transition zone) of about 100 mm width. Fig. 10 shows the TEM and SAED results of the region near the fusion line. The sampling position is shown in Fig. 10(a), which was obtained by FIB. The SAED results indicate the 309L SS substrate near the fusion line has an austenite structure (FCC), as shown in Fig. 10(d), and the element transition zone matrix shows a structure with BCC in Fig. 10(e). Fig. 10(c) displays a TEM image of the element transition zone and granular or rod-shaped precipitates
are found in this region. Fig. 10(f) shows the SAED result, which indicates that those precipitated pellets are mainly metallic carbide. Both M23C6 and M7C3 (M is mainly Fe and Cr) that show an orientation relationship of axis [01-1]//[331] is detected.
Table 3 SEM-EDS results of the chemical compositions (wt. %) of inclusions and pits on 309L SS after immersion in deaerated high temperature water.
Inclusions Pits
Fe
Cr
Ni
Mn
Si
O
S
9.41 22.22
6.50 22.55
0.81 1.92
32.42 7.56
17.86 0.63
27.09 20.48
1.10 0.19
Fig. 6. Optical micrograph of stainless steel near the fusion line.
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Fig. 7. Optical micrographs of (a)309L SS and 308L SS, (b) ferrites in 309L SS, (c) (d) ferrites in 308L SS, (e)(f) SEM maps of ferrites in 309L SS and 308L SS respectively.
3.2. Oxide films formed on the specimen surface after immersion in deaerated high temperature water at 290 C After the immersion test for 168 h, oxide films formed on the weld specimen are characterized. Fig. 11 shows the low
magnification of film on the weld surface. The interface of oxide film formed at various parts is obvious. The region of A508III close to the fusion line shows a bright zone, similar to the decarburization zone. Corrosion pits are detected on 309L SS, as displayed in Fig. 11(c), but obvious pits are not found on the 308L SS surface.
Fig. 8. EBSD results of 309L SS and 308L SS: (a) grain orientation coloring maps (b) grains boundary character maps and (c) inverse pole maps. X indicates the distance from the fusion line.
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Fig. 9. SEM-EDS results (a) Distribution of some elements near the A508III-309L SS fusion line and (b) details.
Fig. 12 shows the SEM images of pits on 309L SS, and the average size of pits is approximately 20 mm in diameter. Fig. 13 shows SEM images of the oxide film formed on the A508III in the weld. Fig. 14 show SEM images of the oxide film formed on stainless steel. Fig. 15 shows the Raman spectroscopy results obtained for the oxide films formed on LAS and SS after immersion in high temperature water, where N1-N6 in the figures indicates the locations far from the fusion line, and the corresponding position (except N1) is shown in Fig. 11. N1, N2 and N3 are located in the bulk of A508III, FGZ and CGZ, respectively, whereas N4 is located in austenite zone, and N5 and N6 are located in 309L and 308L, respectively. The TEM results of the oxide films formed on 309L SS and 308L SS surfaces after immersion in deaerated hightemperature water are shown in Figs. 16e21.
3.2.1. Oxide films formed on A508 III The surface of CGZ is covered with almost no oxide particles, but those on FGZ, ICCGZ and the bulk zone are no different according to the SEM observation, as shown in Fig. 13(c)-(d). The Raman spectra
indicate that the types of the oxide films formed on various regions of A508 III are similar and that the oxide film consists primarily of Fe3O4, as shown in Fig. 15(a). 3.2.2. Oxide films formed on 309L/308L Relatively small and sparse particles are distributed in the fusion zone, as shown in Fig. 14(a). The surfaces of 309L and 308L SS are covered mainly with faceted oxide particles, as shown in Fig. 14(b) and (c). The faceted particles are composed primarily of Fe-rich oxides with low Cr and Ni contents, and the EDS results are shown in Table 4. Considering the low concentrations of oxygen in the oxide film in Table 4, it is thought that the compositions obtained by SEM-EDS were probably influenced by base metal underneath the oxide film. The Raman peaks are characteristics of spinel with AB2O4 on both the 309L SS and 308L SS surfaces, according to the characteristic peaks [24], and the spinel oxides should be (Fe,Ni) (Fe,Cr)2O4 according to EDS results. The oxide particles on 309L SS are denser in distribution than are those on the 308L SS surface.
Fig. 10. (a) Sampling location by FIB for TEM observation, (b) 309L SS substrate near the fusion line, (c) Element transition zone matrix, (d) SAED pattern of stainless steel substrate, (e) SAED pattern of element transition zone e matrix (f) SAED pattern of M23C6 (zone axis [01-1]) and M7C3 (zone axis [331]).
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Fig. 11. Low magnification photos of immersion sample in deaerated water (a) low image panorama (b)interface (c)309L (d)308L.
4. Discussion 4.1. Microstructure
Fig. 12. SEM images of pits on 309L SS after 168 h immersion in deaerated hightemperature water at 290 C.
Fig. 16 show the cross section images of the oxide films on 309L and 308L, showing a thicker inner oxide film on 309L than on 308L. The oxide films show a duplex structure on both 309L SS and 308L SS. The outer layer is composed of large and dispersive oxide particles that are determined to be spinel-type structures by the diffraction pattern (as shown in Figs. 17(I) and 19(I)). The inner layer of the oxide film on 309L SS is 140 ± 20 nm thick (region II in Fig. 17), whereas the inner layer of the oxide film on 308L SS is 60 ± 30 nm thick (region II in Fig. 19). The SAED pattern of region II from the inner layer showed diffraction rings which indicate the presence of nanocrystalline spinel oxides in the inner layer of 309L SS and 308L SS, and the dark area (region III in Figs. 17 and 19) is a matrix with FCC structure, which indicates that the matrix of 309L SS and 308L SS is shows an austenite structure. Figs. 18 and 20 show EDS mappings of the oxide layers formed on 309L SS and 308L SS respectively. The EDS results show that the outer oxide layer on both 309L SS and 308L SS is Fe-rich and that the inner layer is Crrich; further, Ni is enriched in the interface between the oxide and the matrix. The character of the element distribution is revealed in the line scan profiles, as shown in Fig. 21, which agree with the result of EDS mappings of the oxide films.
According to metallographic observations after microhardness tests (shown in Fig. 2), the unstable hardness in the HAZ is due to the black stripe, which shows a higher hardness, and the black stripe should be carbon condensed. The highest hardness near the fusion line may relate to a martensite structure. In fact, increasing carbon and chromium contents favors the formation of martensite [25], and many researchers have found that martensite exists near the fusion zone in dissimilar joints [13,14]. Based on the Schaeffler diagram, martensite will appear at a certain dilution, but in this work martensite structure is not found by TEM, possibly due to sampling position or discontinuous distribution of martensite. A similar primary boundary and secondary boundary relation was also found in other studies [26]. It is worth noting that the distribution of the secondary boundary is discontinuous. Nelson et al. [27] suggested that the primary boundary is a result of epitaxial growth based on the existing substrate during solidification. The formation of the secondary boundary depends mainly on the change in the solidification mode during the solidification process [28]. Molten metal solidification near the fusion line is ferrite (BCC) and then austenite (FCC) solidification in 309L SS, with a change in solidification. Nelson et al. considered the formation of the secondary boundary is to be the result of grain boundary migration during austenite solidification. Yoo et al. [29] studied the effects of thermal aging on the microstructure of Type-II boundaries in Ni-based and low alloy steel dissimilar welds and found that the Type-II boundaries shifted away from the fusion line with an increase in the thermal aging time at 450 C. Therefore, the secondary boundary should be formed by the migration of the austenite boundary. Previous works have found that the secondary boundary, which is parallel to fusion boundary in the filler metal in dilution zone is a potential crack path for SCC [27], which will speed up the failure of the weld. The bulk SS shows a mixed austenite and ferrite structure, and the ferrites in 309L SS and 308L SS show differences in content and shape. From the Schaeffler diagram [30], the ratio of Creq/Nieq in 309L/308L SS falls into the gþd region. Cr is a ferrite-stable element, and a high Cr content is favorable for the formation of ferrite, whereas a higher content of Ni, which is an austenite-stable element in 309L SS, has a negative effect on the stability of ferrite. As shown in Table 5 from SEM-EDS analyses, the ferrites
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Fig. 13. SEM images of oxide film on LAS after 168 h immersion in deaerated high-temperature water at 290 C, (a) film on CGZ (b) film on FGZ (c) film on ICCGZ (d) film on bulk base metal of A508III.
Fig. 14. SEM images of oxide film after 168 h immersion in deaerated high temperature water at 290 C, (a)interface (b)309L (c)308L.
distributed in SS reveal a higher ratio of Cr/Ni than austenite does. Heat input is also crucial to the ferrite content [31], increasing the welding temperature and heat treatment time decreases the content of ferrite. Heat treatment after cladding 309L SS and cladding the second layer of 308L SS are also factors that cause 309L SS to exhibit a higher ferrite content than 308L SS. Relevant research has found that the distribution of ferrite in austenite stainless was related to the ratio of Creq/Nieq and welding speeds which determine the solidification mode of Fe-Cr-Ni alloy [31e33]. A primary ferrite following d transform to g (FA mode) or a primary austenite following g transform to d upon cooling (AF mode) may occur on austenitic stainless steel welds. Chemical composition and welding process conditions determine solidification mode. The following
four solidification modes have been reported by previous researchers [34]: F mode: Creq/Nieq (in wt.%)< 1.25, Lþ g/g
(1)
AF mode: 1.25< Creq/Nieq (in wt.%) < 1.48, Lþ g/Lþ gþd/gþd
(2)
FA mode: 1.48< Creq/Nieq (in wt.%) < 1.95, Lþd/Lþdþg/dþg
(3)
A mode: 1.95< Creq/Nieq (in wt.%), Lþd/d
(4)
where
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Fig. 15. Raman spectra of the oxide films on weld after 168 h immersion in deaerated high-temperature water at 290 C, (a) A508 III (b) 309L/308L SS.
Fig. 16. TEM images of oxide films on (a) 309L SS (b) 308L SS after 168 h immersion in deaerated high-temperature water at 290 C.
Fig. 17. TEM image for area 1 in Fig. 16a) and correspondent electron diffraction patterns of oxide film on 309L SS after exposure in deaerated water. The white dotted line indicates the interface of the outer-layer and inner-layer oxide film.
Creq ¼ Cr þ Mo þ 1.5 Si þ 0.5 Nb
(5)
Nieq ¼ Ni þ 30 C þ 0.5 Mn
(6)
For the present work, the Creq/Nieq of 309L SS is 1.66, and the Creq/Nieq of 308L SS is 1.77. It is suggested that the solidification of
309L SS and 308L SS falls into the FA mode. However, both AF mode and FA mode are observed on 308L SS as shown in Fig. 7(a). As observed in other studies [35e37], FA mode presents a greater amount of ferrite with lathy, whereas AF mode presents skeletal ferrite, similar to the results shown in Fig. 7(c) and (d). Suutala [33] reported that the welding speed and cooling rate have an effect on
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Fig. 18. STEM-HAADF image of oxide film on 309L SS and the corresponding EDS mappings for Fe, Cr, Ni and O after immersion in deaerated water.
Fig. 19. TEM image for area 2 in Fig. 16b) and correspondent electron diffraction patterns of the oxide film on 308L SS after exposure in deaerated water. The white dotted line indicates the interface of the outer-layer and inner-layer oxide film.
Fig. 20. STEM-HAADF image of oxide film on 309L SS and the corresponding EDS mappings for Fe, Cr, Ni and O after immersion in deaerated water.
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Fig. 21. EDS line scan profiles of the oxide film in deaerated water, (a) on 309L SS along the white dotted line a as shown in Fig. 15, (b) on 308L SS along the white dotted line b shown in Fig. 17.
Table 4 SEM-EDS results of the chemical compositions (wt. %) of the oxide particles (shown in Fig. 14) after immersion in deaerated high temperature water.
Particles on 309L SS Particles on 308L SS
Fe
Cr
Ni
O
67.78 60.74
4.74 8.09
2.47 4.39
22.70 22.22
Table 5 SEM-EDS results of austenite and ferrite in 309L SS cladding alloy. Phases
Fe
Cr
Ni
Cr/Ni
Austenite Ferrite
69.20 73.38
17.02 20.61
10.48 3.18
1.62 6.48
the solidification mode. The steady-state dendritic growth indicates the effect of the solidification conditions and a higher steady-state dendritic growth rate favors austenite as the primary phase. A high welding speed and cooling rate are preferred for AF mode. In the weld pool, the temperature gradient and element segregation during solidification are the main factors causing AF and FA mode to occur in the 308L SS surface layer.
4.2. Effects of LAS microstructure on the oxide film formed in deaerated high-temperature water Fe3O4 particles are detected on the LAS surface film in deaerated water, as shown in Fig. 15(a). From the potential-pH diagram [38,39], Fe3O4 is more stable at low potential whereas Fe2O3 is more stable at high potential, which suggests that Fe3O4 is easier to form in deaerated high-temperature water. Meisel [40] predicted the iron oxide or iron hydroxide type by calculating the effective Gibbs free energy Geff, where the generated Geff of the nonstoichiometric compound Fe3O4 is the minimum value. Fe3O4 should be the product of LAS corrosion in deaerated high-temperature water. The Raman spectra show that there is no significant difference in oxide species on each region of the LAS surface, whereas the surface oxide distribution on CGZ is quite different from that on FGZ, ICCGZ and the bulk metal, as displayed in Fig. 13. The FGZ consists mainly of ferrite coinciding and has a lower carbon content than do other regions of LAS due to decarburization. Liu et al. [41] studied the influence of carbon content on corrosion resistance of low alloy steel and found that the corrosion potential of corrosion layer becomes more positive, the corrosion current density
decreases and the oxide film has more protection performance with a decreasing carbon content of LAS. Clover et al. [42] considered that the main reason why the pearlite phase is worse than ferrite in corrosion resistance is that pearlite is a eutectoid body of ferrite and cementite; further, as the potential of the two-phase is different, there is severe micro galvanic corrosion in pearlite phase and cementite because the cathode will accelerate the corrosion of ferrite in pearlite. Therefore decarburization may be the main reason for the CGZ showing almost no particles on the film compared to other regions. 4.3. Effects of stainless steel microstructure on the oxide film formed in deaerated high-temperature water The oxygen concentration has a great effect on the electrochemical corrosion potential of metal in high-temperature water [43]. Potential-pH diagrams predicted that the spinel-type oxide is more stable on Fe-Cr-Ni alloy in low-oxygen concentration water. Stellwag [44] suggested that the outer oxide film is formed by the dissolution of metal ions in deposition and that the rate of Fe atoms migration is faster than those of Ni and Cr in oxide. Thus, the outer oxide particles are Fe-enriched, whereas Cr and Ni generally remain in the inner layer. The oxide films formed on 309L SS and 308L SS were observed by SEM and TEM. Combined with the EDS results, the surface oxide particles are Fe-rich spinel, consistent with the Raman spectroscopy results. Kuang et al. [45] studied the effect of the dissolution of Ni on the oxidation of 304 SS at 290 C and found that spinel particles started to appear when low Ni2þ concentration transferred to high in solution, which suggests that Ni is favorable for spinel nucleation. 309L SS matrix has a higher Ni content. Therefore, the spinel-type particles distributed on 309L SS are denser than those distributed on 308L SS surfaces. The oxidation rate of Ni is lower than those of Fe and Cr in hightemperature water. Because of its low oxygen affinity, Ni is the least likely to be oxidized compared with Fe and Cr [46,47]. Researchers have found Fe and Cr atoms tend to deviate from the surface more easily than do Ni atoms and that Ni enrichment appeared near the inner oxide film in studies on stress corrosion behavior of Fe-Cr-Ni alloy that used the molecular dynamics model [22]. Therefore, the migration rate of Ni into the oxide film is very low, resulting in an enrichment of Ni in the interface of the inner layer and matrix, which can be clearly observed in the EDS mapping. The TEM results show that the oxide inner layer on 309L SS is thicker than that on 308L SS. EDS mapping and line scan profiles
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both reveal that the inner oxide layer is consisted of Cr-rich oxides. The enrichment of Cr in the inner layer contributed to the corrosion resistance. The oxidation rate constant of Ni-based alloys in simulated PWR primary water was calculated [17,18,48], and the oxidation rate constant of alloys 600 and 182 with 15% Cr content was higher than that of alloy 690 with 30% Cr content. Das et al. [22] found that atoms diffused slower in Cr-rich oxide than in Fe-rich and Ni-rich oxide and that a Cr-rich oxide layer is counterproductive to the growth of oxide film. The inner layer of the oxide film formed on 316L SS with a higher Cr content was thinner than that on 9Cr-26Ni-1.5Mo-5Si alloy as reported in Ref. [45]. The oxidation rate of stainless steel in high-temperature water is mainly determined by the Cr-rich inner oxide layer. Data from multiple measurements indicate that the content of Cr of the inner oxide layer on 309L SS is 24.4% (wt%) whereas the content of Cr in the inner oxide layer on 308L SS is 33.47% (wt%). The growth of oxide film depends on the metal atoms migrating from matrix to oxide film and oxygen diffusing into oxide film. The oxide film with a higher Cr content can mitigate more effectively the transport of metal atoms and oxygen atoms, which is supposed to result in a lower oxide film. However, experimental data show that Cr content of 309L SS substrate is higher than that of 308L SS substrate. The higher Cr content in the inner oxide layer of 309L SS can be related to the effect of ferrite on oxidation. Cao et al. [23] found that higher Cr content led to a higher ferrite content and reduced the amount of oxidized metal. Post weld heat treatment (PWHT) had no significant effect on the formation of oxide film. Their work did not emphasize the effect of ferrite on the oxide formation. Previous works have reported that the interaction between the austenite and the ferrite would affect the corrosion potential of dual-phase steel [49]. The galvanic interaction influence of a-phase in stainless steel has been investigated [50,51]. a-phase could move the open circuit potential of g-phase toward the noble direction, and the passive film impedance of single g-phase increased after the contact of the two phases. The oxide film on the duplex stainless steel showed better corrosion resistance. Lu et al. [52] reported that the crack growth rates in the weld metal with ~6.5 wt% ferrite content were higher than those in the weld metal with ~8.5 wt% ferrite content in dissolved oxygen water at 288 C. Those studies indicated that the ferrite content of SS affects the property of oxide film formed on SS. Compared with 309L SS, 308L SS has a higher ferrite content but a lower Cr content. Ferrites in stainless steel may improve the Cr content of oxide film, which indicates a lower growth rate of oxide film on 308L SS than on 309L SS. However, the detailed reasons for the effect of ferrite to Cr concentration in oxide film require further study. Many corrosion pits could be found only on the 309L SS surface after immersion. Obviously inclusions are observed in 309L SS but are not observed in 308L SS before the immersion test. Through observation and statistics, the distribution and size of inclusions and pits were found to be similar. The EDS results of inclusions and pits are displayed in Table 3. The Si and Mn contents decreased sharply in pits compared with inclusions, which suggest inclusions may have dissolved in deaerated high-temperature water and formed the pits. The preferential dissolution of Si in oxide films in high-temperature water has been reported [53]. This work suggests that silicon oxide has high solubility in high-temperature water which agrees with this result. It confirms that corrosion pits on the film of 309L SS surfaces are caused by inclusion dissolution. These corrosion pits will act as channels for crack propagation and accelerate the stress corrosion cracking. Nonmetallic inclusions in steels affect corrosion resistance, such as pitting corrosion, stress corrosion cracking and hydrogen-induced cracking [54,55] resistance.
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5. Conclusions The microstructure and properties of the oxide films formed on various locations of a A508III/309L/308L weld clad in simulated deaerated PWR high-temperature water are characterized as follows. 1. The HAZ in the LAS side consisted of CGZ, FGZ and ICCGZ. A high hardness zone in the SS zone near the fusion line was identified, and attributed to the C-enrichment. M23C6 and M7C3 precipitates in element transition zone were identified by TEM. 2. Ferrite content was approximately 9% in 309L SS and approximately 12% in 308L SS. Both FA and AF mode were found in 308L SS, but only FA mode was found in 309L SS cladding. 3. The surface oxide films on LAS HAZ and base metal in deaerated high-temperature water were mainly Fe3O4 type, where CGZ exhibited less oxide coverage than did other regions in the LAS HAZ. The oxide films on stainless steel cladding surface were mainly spinel oxides. More corrosion pits appeared on 309L SS than on 308L SS after immersion in deaerated high-temperature water due to the dissolution of inclusions. 4. The oxide film formed on SS cladding in high temperature water can be divided into two layers: the Fe-rich spinel outer oxide layer and the Cr-rich oxide inner layer. Ni-enrichment at the oxide/substrate interface is observed. The oxide film formed on 309L is thicker than that on the 308L, which is thought to be related to the effect of ferrite on oxidation. Acknowledgments This work has been supported by Natural Science Foundation of China (NSFC No. 51771107), National Key R&D Program of China (No. 2017YFB0703002), and Shanghai Municipal Commission of Economy and Informatization (No. T-221715003). The support from the Instrument Analytical and Research Center, Shanghai University is acknowledged. References [1] [2] [3] [4] [5]
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