martensitic steels in flowing supercritical water

Journal Pre-proof Oxidation Behavior of Ferritic/Martensitic Steels in Flowing Supercritical Water Quanqiang Shi, Wei Yan, Yanfen Li, Naiqiang Zhang, Yiyin Shan, Ke Yang, Hiroaki Abe

PII:

S1005-0302(20)30009-8

DOI:

https://doi.org/10.1016/j.jmst.2020.01.009

Reference:

JMST 1876

To appear in:

Journal of Materials Science & Technology

Received Date:

18 April 2019

Revised Date:

28 May 2019

Accepted Date:

3 July 2019

Please cite this article as: Shi Q, Yan W, Li Y, Zhang N, Shan Y, Yang K, Abe H, Oxidation Behavior of Ferritic/Martensitic Steels in Flowing Supercritical Water, Journal of Materials Science and amp; Technology (2020), doi: https://doi.org/10.1016/j.jmst.2020.01.009

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Research Article

Oxidation

Behavior

of

Ferritic/Martensitic

Steels

in

Flowing

Supercritical Water Quanqiang Shi1, Wei Yan1, Yanfen Li1, Naiqiang Zhang2, *, Yiyin Shan1, KeYang3, *, Hiroaki Abe4 1

Key laboratory of nuclear materials and safety assessment, Institute of Metal Research, Chinese

Academy of Sciences, Shenyang 110016, China 2

Key Laboratory of Condition Monitoring and Control for Power Plant Equipment of Ministry of

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Education, North China Electric Power University, Beijing 102206, China Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China

4

School of Engineering, The University of Tokyo, Tokai, Ibaraki, 319-1188, Japan

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3

* Corresponding authors.

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E-mail addresses: [email protected] (K.Yang); [email protected] (N.Q. Zhang).

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[Received 18 April 2019; Received in revised form 28 May 2019; Accepted 3 July

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2019]

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The oxidation behavior of two Ferritic/Martensitic (F/M) steels including novel SIMP steel and commercial P91 steel were investigated by exposure to flowing deaerated

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supercritical water (SCW) at 700 °C for up to 1000 h. The kinetic weight gain curves follow parabolic and near-cubic rate equations for SIMP and P91 steels, respectively. XRay Diffraction analysis showed the presence of magnetite and a spinel phase in flowing SCW for both steels. The morphology and structure of the oxide scales formed on these two steels were analyzed. The relationship between the microstructure and oxidation behavior and the reason that SIMP steel showed better oxidation resistance than P91 steel were discussed.

Key words: Ferritic/martensitic steel, Supercritical water, Oxidation, Oxide scale, SIMP steel

1.

Introduction The supercritical water-cooled reactor (SCWR) is considered to be one of the most promising

candidates of Generation IV advanced nuclear reactor concepts due to its high thermal efficiency and

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simple design without steam generators and steam separators [1-3]. The thermal efficiency of the SCWR is expected to be 44%, compared to 33% for the current generation of light water reactors (LWRs) as the operation conditions such as outlet temperature and pressure increase from ~285-320 oC / ~7-15 MPa to ~550 oC / ~25 MPa [1, 4]. However, supercritical water (SCW), at a temperature and pressure [5, 6]

, is severely corrosive to

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above the thermodynamic critical point of water (374 oC / 22.1 MPa)

structural materials. Thus, the oxidation behavior of materials at SCW conditions has been one of the

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main issues under investigation [7, 8].

9%-12% Cr F/M steels have been identified as candidates for fuel cladding and core structural

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components in some Generation IV SCWR concepts. F/M alloys have already been put into application in the energy industry because of their better high temperature strength and creep resistance, higher thermal conductivity, lower swelling and activation under irradiation, lower cost, lower thermal

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expansion coefficient and better resistance to inter-granular stress corrosion crack (IGSCC) than austenitic stainless steels [9-16]. Oxidation has been observed as the predominant corrosion phenomena

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of 9%-12% Cr F/M steels exposed to environments such as high temperature vapor/steam [17] and SCW. In general, steels with high Cr content show superior oxidation resistance compared to those with low [18]

. The natures and structures of the oxides formed on 9% Cr F/M steel P91 exposed to

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Cr content

deaerated SCW are similar to those obtained in wet air at the same temperature

[19]

. To date,

international programs investigated on oxidation behavior in SCW were focused on representative commercial 9%-12% Cr F/M steels, such as P92 [20], HCM12A, NF616 [21] and HT9 [22]. SIMP steel is a novel modified 9%-12% Cr F/M steel which has been developed by our research team for application in the harsh environment caused by the liquid LBE interaction with the containment at high temperature [23-28]. The test results conducted on one-ton and five-ton scales SIMP steels showed that

SIMP steel has more potential as a candidate structural material for the spallation target in Accelerate Driven Sub-critical (ADS) system as its mechanical properties

[25, 27]

, high temperature oxidation

resistance in air [23, 28] and oxidation resistance in static liquid LBE saturated by oxygen [24] was better than that of T/P91. However, research on oxidation behavior of SIMP steel in SCW environment is seldom reported. Therefore, the present work aimed to study the corrosion behavior of SIMP steel at 700 oC in SCW, taking P91 steel as a comparison. Based on the research results such as the microstructure of the oxide scale and the analysis of the reason for superior high temperature oxidation resistance of SIMP

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steel, a kinetic model was proposed to describe the nucleation, growth and degradation of oxide scale formed on the surface of SIMP steel in SCW. Experimental

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2.

Two 9%-12% Cr F/M steels including the SIMP and commercial P91 steel were studied in this

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work. The chemical compositions of the two steels are listed at Table 1. The SIMP steel was normalized at 1050 oC for 0.5 h, air cooled, then tempered at 760 oC for another 1.5 h before final air cooling, and

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the P91 steel was normalized at 1050 oC for 20 min, air cooled, and then tempered for 1 h at 780 oC before final air cooling. Through the above heat treatment, both steels obtained a martensitic lath microstructure, with Cr23C6 carbides along the lath and the prior austenite grain boundary [14]. There

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were a larger number of carbides in SIMP than that of P91 steel [23, 24]. After heat treatment, specimens with dimension of 20 mm × 10 mm × 2 mm were cut from the

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bulk of SIMP and P91 steel by electric discharge so that the long axis of the specimens was in the rolling direction and each specimen had a small hole for attaching it in the autoclave. Prior to the

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corrosion tests, the samples were ground and surface polished up to 2000 grit by following the standard metallographic procedures. Oxidation tests in flowing SCW at 700 °C were performed in a continuous flowing SCW

experimental facility at a pressure of 25 MPa with a Dissolved Oxygen (DO) content of 2 ppm and a water flow rate of 5 mL/min through a 5 cm2 area cross-section. The schematics of the SCW experimental facility can be seen in Reference [8]. Ultrapure water with an electrical conductivity less than 0.1 μS/cm was obtained using an ion exchanger. The water was deaerated by heating to 100 °C

and bubbling with pure nitrogen gas. The DO content in the reaction system was controlled by injecting mixed argon and oxygen gas, and it was corrected by a mass flow controller according to the DO sensor behind the backpressure valve. The flow rate and pressure were controlled through the highpressure metering pump and the backpressure valve. Samples were placed in the autoclave using platinum wires and ceramic insulators to avoid galvanic effects. The exposure periods were 200, 400, 600, 800 and 1000 h. In each test, specimens were removed from the equipment after reaching their designed exposure time. The specimens with longer exposure times were cooled to room temperature and reheated to the designed testing temperature to complete

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the rest experimental time. After the tests, the specimens were weighted using a METTLER TOLEDO balance with a sensitivity of 0.01 mg. The oxides formed on the surface of the samples were identified by means of X-Ray Diffraction (XRD), and the high definition scans were performed at a rate of 0.02o/s in the 2θ interval ranging from 15o to 85o. The surface morphologies of the oxides on the samples were

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observed under Scanning Electron Microscopy (SEM), the specimens were coated with a thin layer of gold to allow charge dissipation during the analysis. The cross section of the oxides on the samples

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were observed under SEM and Electron Probe Micro-Analysis (EPMA) after grounded and polished by following the standard metallographic procedures until a mirror finish surface was produced. To

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improve the image quality, the samples were also coated with a thin layer of gold. Energy dispersive

3.

Results

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3.1 Weight gain

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X-ray (EDX) analysis was conducted on the samples when required.

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The relationship between the weight gain and the exposure time can be expressed as [29]: ∆W = ko exp (-E / RT) tn = kp tn

(1)

where ∆W represents the weight gain of the steels per unit area (mg/dm2), E is the activation energy for oxidation, R is the ideal gas constant (8.314 J/(K mol), ko and kp are the oxidation rate constants, T is the temperature in Kelvin, t is the exposure time, and n is an exponent describing the time dependence of the oxide growth. The weight gains as a function of time for SIMP and P91 steel

exposed to deaerated flowing SCW at 700 °C under 25 MPa are shown in Fig. 1. The parameter of R2 standards for measuring the significance of the n of the fitting line differing from zero

[21, 30]

. Fig. 1

also shows a plot of los weight gain versus los time. The weight gain increased as the exposure time increased, but the weight gain of SIMP steel was almost half of that of P91 steel. The time exponents obtained from the experimental results at 700 °C are 0.47786 for SIMP steel and 0.4011 for P91 steel, which indicated that the weight gain of the SIMP steel obeyed a parabolic rate equation and that of P91 is between the cubic and parabolic rate equation. The oxidation rate of SIMP steel was slower than P91 steel, which implied that the higher Cr and Si

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content in SIMP steel resulted in the formation of more protective oxide layer and was consistent with the previous research [23, 24]. The correlation coefficients (R2) of SIMP and P91 steels were more than 0.99, which indicated that the fitting curves had a good fitting quality for the curves.

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3.2 XRD analysis

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The XRD analyses of the samples also provide important information on the oxidation phenomena. Phase analyses of the surface oxide formed on SIMP and T91 steels by XRD after

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oxidation exposed to deaerated flowing SCW at 700 °C under 25 MPa for different time are shown in Fig. 2. The XRD results obtained from both alloys were similar and the octahedral magnetite (Fe3O4) was the only detected phase. However, the orientation of the magnetite (Fe3O4) formed on the two

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steels were totally different, the orientation of the magnetite formed on SIMP steel was mainly the (311) plane, indicating a preferred oxide growth direction in this case, while the orientation of the

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magnetite formed on P91 steel was along the (400) plane.

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3.3 Oxide morphology

Fig. 3 shows the surface morphology images of the oxide scales on SIMP and P91 specimens as

a function of exposure time after being exposed to flowing deaerated SCW at 700 °C under 25 MPa. As shown in Fig. 3, no oxide exfoliation was observed over the duration of experimental procedure. The grain size of the magnetite, identified and confirmed by XRD and EDS analysis, gradually grew up as the exposure time increased, and the average size of magnetite formed on SIMP steel was a bit larger than that of P91 steel and much smaller grain size of magnetite located at the surface of P91

steel during the entire exposure time. Another distinct characteristic was that most of the magnetite crystals on the SIMP steel presented just one triangular facet parallel to the surface while the magnetite on the P91 steel mostly presented tetrahedral facet, this different orientation was reflected in the XRD analysis. Micro and large cracks both were also observed in the oxide scales formed on the specimens exposed to 700 °C for 400 h, which was about 200 h ahead of that reported by Zhang et al [8]. As the SCW temperature increased from 600 °C to 700 °C, the oxidation rate was significantly enhancing [8, 29]

, the growing stress in the oxide scales accumulated and caused the crack initially formed to

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propagate through the oxide scales, which indeed increased the effective diffusivity with the exposure time. It was interesting to note that the surface oxide of SIMP specimens presented ridge-like morphologies especially at the early stage of oxidation, as shown in Fig. 4. The size of the magnetite oxide at the bottom of the ridge was much smaller than at the top of the ridge. The phenomenon of

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ridge-like oxide indicated that the oxidation procedure of SIMP steel was not even, which was proved by the cross section of oxide structure, and has been interpreted in detail in the discussion section.

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The adhesion between the oxide scale and SIMP steel matrix exposed to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h was poor and the oxide layer peeled off by the conductive tapes

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used during surface oxide morphology observation. The enlarged image and EDS analysis of the exfoliated oxide are presented by Fig. 5 and Table 2. As shown in Fig. 5, the morphologies and sizes

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of oxides located at A, B, C were different. The oxides at the A and B positions were both magnetite confirmed by EDS and XRD analysis. The grain size of magnetite at A position was 25-30 µm and cracks were observed, however, the grain size of magnetite at B position was 2-5 µm, just one sixth of

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position A, and it did not seem compact as there were pores between the magnetite particles. Location C is assumed to be the original interface between the SCW and SIMP steel as the scratches left during

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the polishing process could be observed (Fig. 5(C)). The oxides, which exhibited lamellar-like structure and were different from Fe3O4, were a Fe-Cr spinel with Mn enrichment as confirmed by EDS analysis.

3.4 Oxide structure The oxide films generally consisted of double layers: an outer layer and an inner layer with seldom

internal oxidation zone (IOZ) at the oxide/metal interface for both steels. Fig. 6 shows SEM-BSE images and EDS element line analysis of the oxide scale structure formed on SIMP and P91 steels after exposure in flowing deaerated SCW at 700 °C under 25 MPa for 1000 h. The outer layer of SIMP steel was characterized by a rougher outer surface than P91 steel and a band of pores, with size of 10-15 µm, presented between outer and inner layer interface. There was considerable porosity located at the entire inner layer and the interface between the oxide scale and matrix was not even while the inner layer of P91 steel was much denser and the oxide/substrate interface was much smoother. Cracks in the oxide of P91 steel penetrated throughout the oxide scale.

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Fig. 7 shows SEM-BSE images and EDS element line analysis of the oxide scale structure formed on SIMP steel after exposure in flowing deaerated SCW at 700 °C under 25 MPa for 200 h. The chromia layer was formed due to the higher oxygen affinity of Cr and its relatively rapid diffusion within the grain boundary [3, 31, 32], resulting in a magnetite layer that was thinner than other locations

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as the diffusion of Fe in chromia layer was rather slow. An IOZ consisting of some small oxide precipitates formed along the grain boundaries (the dark area shown in Fig. 7) was observed, which

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indicated that the lath and grain boundaries were oxidized first and the oxygen penetrated into the metal matrix along the lath and grain boundaries [20, 33, 34].

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EDS was used to determine the composition of the oxide films. The composition profiles of the SIMP and P91 steels exhibited two distinct oxide/metal ratios that correspond to two different oxides.

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The oxygen content of the distinct oxide layers was similar: the outer oxide was predominantly iron oxide while the inner layer contained a significant amount of chromium, with silicon mainly enriched near the matrix of SIMP steel, and the oxygen potential dramatically decreased at the Fe-Cr spinel and

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matrix interface. The outer oxide layer of both steels had an oxygen-to-metal (O/M) ratio ~1.3, consistent with magnetite, as identified by XRD analysis. The O/M ratio of inner layer was 1.2-1.3

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and was consistent with an iron-chromium oxide (Fe, Cr)3O4. The grey zone at the interface of the FeCr spinel and matrix of SIMP steel was silicon enriched oxide, confirmed by EDS analysis. The element segregation in the Fe-Cr spinel of SIMP steel was more severe than that of P91 steel, with variations of the Cr and Fe concentrations, and silicon was mainly enriched at the interface of the matrix and oxide scale. The SEM-BSE images and the oxide scale thickness including the outer over inner layer ratio of SIMP and P91 steels in flowing deaerated SCW at 700 °C under 25 MPa for different time are shown

in Fig. 8 and Fig. 9. The oxide thicknesses of SIMP steel were estimated to be about 41.23, 56.69, 67.05, 78.05 and 89.31 µm and the thickness of outer oxide layer was about 21.02, 28.34, 32.66, 37.40 and 42.31 µm for exposure time of 200, 400, 600, 800 and 1000 h, respectively. The oxide thicknesses of P91 steel were determined as 71.33, 101.94, 110.54, 127.81 and 143.22 µm and the thickness of outer oxide layer was about 37.68, 54.30, 59.60, 70.24 and 79.85 µm at exposure time of 200, 400, 600, 800 and 1000 h, respectively. The oxide thickness of SIMP and P91 steels increased as the corrosion time increased, and the thickness of the oxide scale of SIMP steel was rather smaller than that of P91 steel at the same corrosion time. The oxide thickness was found to be proportional to the

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weight gain. The total thickness showed a similar behavior as the weight gain and the index of the fitting line of the oxide scale thickness was consistent with weight gain curve of the SIMP and P91 steels.

The most pronounced difference between the SIMP and P91 steels was the pore positions in the

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outer magnetite layer of the scales. At first, the pore position of SIMP steel was at the interface of the outer and inner oxide layer, but as the exposure time increased and magnetite grains grew, the pores

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were generated beneath the outmost magnetite layer and pore band ultimately formed at the oxide scale and matrix interface after exposure to SCW for 1000 h. However, the pore location on the P91 steel

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was just beneath the outmost magnetite layer and did not move significantly during the entire exposure period. No oxide exfoliation was observed during the entire exposure period, only the formation of

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cracks penetrating almost perpendicularly through the oxide scales on P91 steel. The thickness ratios of the outer to inner oxide layers of SIMP and P91 steels were between 0.90 and 1.26. However, the trend of the ratio variation over the whole exposure period was different for the two alloys. The ratio

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of SIMP steel decreased from 1.04 to 0.90 while the ratio of P91 steel increased from 1.12 to 1.26. Note that SIMP steel consistently had the lower ratio (thicker inner oxide) and displayed the lower

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oxidation rate. In contrast, P91 steel exhibited the higher ratio and higher weight gain. The elemental contents of the inner Fe-Cr spinel layer of SIMP and P91 steels taken from Fig. 8

obtained from EDS analysis is shown in Table 3. The chromium and silicon contents in the oxide layer were almost twice as much as the content of the alloys and remained the same during the entire exposure time. The chromium and silicon content of SIMP steel was more than that of the P91 steel, in particular the silicon content in SIMP steel was almost seven times greater than that in P91 steel. The outer magnetite layer was composed of large columnar grains and had a simple structure

while the inner Fe-Cr spinel was much more complex due to the presence of various types of oxides located in various positions in the oxide layer. The inner Fe-Cr spinel layer was likely to be more protective since its chromium content was much higher than that of the base metal, while the outer layer was thought to be non-protective [35]. Fig. 10 shows the EPMA images of the Fe-Cr spinel and diffusion layers on SIMP and P91 steels after exposure in flowing deaerated SCW at 700 °C under 25 MPa for 1000 h. The chromium, accompanied by silicon, in the Fe-Cr spinel of SIMP steel was segregated to form an iron-chromium elemental separation. The Fe-Cr spinel on P91 steel could be divided into two distinct (A and B) regions. In region A, which was closest to the P91 matrix,

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chromium precipitates accompanied by silicon were segregated intermittently while the chromium and silicon were distributed evenly in zone B beneath the magnetite layer, which was not obvious as compared to SIMP steel. Besides the regular laminated structure, which was parallel to the oxide/metal interface, in the Fe-Cr spinel layer near the P91 substrate was observed, while the Fe-Cr spinel structure

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of SIMP steel exhibited netlike with the depletion of Fe. The EDS line scanning results and EPMA mapping illustrated that the chromium enriched oxide was distributed as striation within the oxide

steels by the previous tests [20, 29, 33, 36-38].

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scale parallel to the oxide/metal interface, which had also been observed in the 9% Cr and 11% Cr F/M

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The STEM images and element scanning of the interface between the oxide scale and matrix of SIMP steel are shown in Fig. 11. A continuous and compact silica layer with a thickness of about 400

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nm, was formed after exposure in flowing deaerated SCW at 700 °C under 25 MPa for 1000 h, which was consistent with the observation of SEM images in Figs. 6 and 7. There were some silica sites

[3, 33]

Discussion

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4.

.

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located at the inner Fe-Cr spinel layer, which was porous due to the outward diffusion of iron cations

4.1 Oxidation resistance difference Based on the above experimental results and analysis, the oxide morphologies and structures formed on both steels exposed to SCW consist of magnetite and a Cr-rich Fe-Cr spinel with no obvious IOZ at the oxide/metal interface, which is consistent with previous research work focused on the corrosion behavior and mechanisms of F/M steels in SCW and other environments

[16, 29, 39-45]

. The

SCW oxidation resistance of SIMP steel was better than that of P91 steel as the weight gain of SIMP steel was much smaller than that of P91 steel and was consistent with the oxide scale thickness. Since the outer magnetite layer was thought to be non-protective [35], the analyses of the inner Fe-Cr spinel and interface between the oxide scale and matrix focused on identifying the reason that SIMP steel showed better oxidation resistance than P91 steel. The columnar grain Fe3O4 outer layer was formed at the SCW/oxide interface due to the outward diffusion of Fe cations, and the inner Fe-Cr spinel layer was formed at the oxide/metal interface when the oxygen penetrated to the metal matrix through pores and grain boundaries. The IOZ was formed [16, 24, 29, 39-45]

.

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because of oxygen diffusion ahead of the oxide along the lath and grain boundaries

Hematite was not formed at the outmost surface in these tests because the SCW was deaerated [29]. The rate controlling process was the outward diffusion of Fe cations since if oxygen diffusion was the rate controlling step then the IOZ would be unlikely to formed as the oxygen would not diffuse ahead of [8, 39, 41-44]

. The main difference between the SIMP and P91 steels was the chemical

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the oxide

composition, which had a significant effect on the corrosion behavior of F/M steels in SCW [44]. The

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distinct differences in the Fe-Cr spinel and diffusion layers between the SIMP and P91 steels were the porosity distribution left by Fe outward diffusion, and the interface irregularity of Fe-Cr spinel (Figs.

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8 and 10), which was attributed to the differences in the chromium and silicon content and microstructure. The chromium and silicon concentration of the Fe-Cr spinel layer of SIMP steel was

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more than that of T91 steel (Table 3), in particular, the silicon content of the SIMP steel was almost seven times as much as that of P91 steel. Chromium and silicon were so stable in the octahedral sites of Fe-Cr spinel compared to iron

[39]

and its presence in the spinel structure inhibited Fe cations

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diffusion, thus slowing down the diffusion of iron. The higher the chromium and silicon contents in the Fe-Cr spinel, the lower the iron diffuse rate

[8, 46-49]

. The addition of 0.52 wt.% Mn in the SIMP

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steel led to it being enriched at the outmost of inner Fe-Cr spinel layer (Fig. 5 and Table 2) because the diffusivity of Mn through the initial chromia is approximately two order of magnitude faster than that of Cr [50].

The carbides of the SIMP steel existed at the lath and boundaries, linked together into chain-like structures [23, 24] and were oxidized by the oxygen that diffused ahead of the oxide along the lath and grain boundaries and formed a chromium and silicon-rich oxide mesh-like diffusion layer due to the thermodynamic stabilities of the oxide [46], impeding the iron migration outwards to form magnetite.

A continuous and compact silica layer was formed at the oxide/metal interface of the SIMP steel after exposure in SCW up to 1000 h (Fig. 11), which might have increased the corrosion resistance by providing a barrier to outward diffusion of Fe cations [46]. The above two points were the reason that the corrosion resistance of SIMP in SCW was much better than P91. The more irregularity of the oxide/substrate interface of SIMP steel was attributed to the higher Cr content than P91 steel, which could form a more protective chromia layer at the initial stage of oxidation due to the different diffusion rates of Cr along the grain boundaries [49, 51]. The inner Fe-Cr spinel layer was preferentially developed along the grain boundaries due to their lower negative

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formation energy and the higher Cr diffusion rates along the grain boundaries [52] (Fig. 7). However, the chromia layer could not protect the matrix as the oxidation time increased up to 1000 h because the chromium content of SIMP steel was less than the critical chromium concentration [46, 53, 54] required to transform a less protective Fe-Cr spinel into more protective external oxide (e.g. Cr2O3). Therefore,

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some regions would be Cr depleted near the initial chromia layer which led to a lack of protection of the base metal, and Fe diffused outward along grain boundaries to form Fe3O4

[20, 40, 55]

. As a more

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compact Fe-Cr spinel and silica layer formed in SIMP steel exposed to SCW up to 1000 h, the outward diffusion rate of Fe cations was slower compared to P91 steel. However, the grain size of magnetite

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was still increasing and the Fe cations were insufficient to meet the needs of magnetite growing of the SIMP steel, so the porosity accumulated beneath the magnetite and a pore band (Fig 6 and Fig. 7) was

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formed at the outer/inner oxide layer interface gradually through the outmost magnetite swallowing up the small size magnetite beneath [39] as the exposure time increased up to 1000 h. The ratio of the thickness of the outer magnetite Fe3O4 layer to the inner Fe-Cr spinel layer could

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be interpreted by the “available space model” [41-43]. At the beginning, the ratio of SIMP and P91 steels was almost the same and equal to 1.1, which was calculated by the “available space model”

[42]

.

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However，the ratio for P91 steel increased by 14.5% while the ratio of SIMP steel was decreased by 18.2% as the exposure to SCW increased to 1000 h (Fig. 9). As mentioned in the above corrosion mechanism of F/M steels, the outward diffusion flux of Fe cations from the metal to SCW/magnetite interface led to vacancy formation in the matrix. Vacancy accumulation could generate formation of cavities, which was the new formed site of Fe-Cr spinel. The cavities were partly annihilated due to the lower Cr content of P91 steel and the distribution of Cr was more even near the magnetite/Fe-Cr spinel interface (Fig. 10) by high compressive stress in the oxide during its growth

[56, 57]

, which

resulted in thickness loss of Fe-Cr spinel layer. On the other hand, because of the presence of more Cr and Si in SIMP steel

[58-60]

, the accumulated cavities were almost transformed into Fe-Cr spinel and

very few cavities were annihilated at the oxide/metal interface[61, 62]. The pores bands formed gradually implying the vacancies, which were located in the magnetite, was annihilated in the magnetite grain and accumulated at the outer/inner layer interface resulting in the thickness loss of the outer magnetite layer. So, the ratios of the two F/M steels exhibited inverse deviation from the expected value. Cracks in the oxide scale of both steels were observed after exposure at 700 oC deaerated SCW. The cracks probably resulted from the thermal expansion mismatch between the oxide scale and the

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alloy matrix on cooling during the experimental procedure. The volume change of magnetite calculated after cooling from 700 oC is about 1.78%, which is twice more than that for the alloy, 0.81%. (Volume expansion coefficient α for magnetite is 20.6 × 10-6 /K (293-843K) and 50.1 × 10-6 /K (843-1273 K) [63]

, and F-M steel was 11.9 × 10-6 /K [64]). The tensile stress generated in the oxide scale will be high

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enough to create micro-cracks because the thickness of the oxide scale on the samples of this

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4.2 Kinetic model of oxide layer growth

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experiment was very high.

A kinetic model of oxide scale growth is described below according to the experimental results, as shown in Fig. 12. When the material was exposed to a SCW environment, chromia layer was formed

[8, 29, 46]

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at the initial stage of oxidation based on the higher oxygen affinity and thermodynamic stability of Cr at the grain boundaries because of the higher diffusion rate at grain boundaries than grain [49, 51]

. However, the low Cr content of SIMP steel lead to Cr depletion and Fe

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compared to Fe

enrichment in the grain boundaries near the surface, and the discontinuous chromia layer is not

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sufficiently protective, leading to a high rate of oxidation far from the initial chromia layer and formation of a thicker surface oxide layer due to the Fe diffused outward along the grain boundaries and caused short-circuit paths to form magnetite Fe3O4 layer

[20, 40, 55]

, as shown in Stage B. The

porosity was attributed to the formation of an Fe3O4 phase with vacancies at interstitial and octahedral sub-lattices to account for deviations from stoichiometry [39, 54, 65, 66], and the vacancies collapsed into pores when the vacancy concentration in Fe3O4 was high enough [3]. As the exposure time increased, the iron cations diffused from the oxide/metal interface to the

outer oxide layer and reacted with oxygen to form a magnetite layer. At the same time, oxygen anions diffusion to inner oxide/metal interface to form the chromium-rich oxide, and then Cr2O3 further reacted with Fe and formed Fe-Cr spinel

[20, 41-43, 54]

. The Cr diffused outward through the grain

boundaries, resulting in the Cr content in the oxide/metal interface, where the Cr rich oxide could develop and Cr depleted beneath the substrate again [36]. The formation of driving force of Fe-Cr spinel was reduced due to the longer diffusion paths of Fe cations, Cr depletion and the lower oxygen potential in the deeper grain boundaries at the internal oxidation front tip [48], as shown in Stage C. The silicon content of the SIMP steel was sufficient to form continuous and compact silica layer

than Cr

[67]

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at the oxide/metal interface but it would take a long time as the diffusion rate of Si was rather lower . during the initial corrosion period, the silicon in SIMP steel might first be oxidized into

discrete SiO2 particles due to the more thermodynamics stability of silica than Cr2O3 [46] and acted as nucleation sites for Cr2O3, so that the formation of transient oxides was suppressed and the overall [61, 68, 69]

. In addition, the Si increased the effective inter-diffusion

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oxidation resistance was enhanced

coefficient of Cr in the metal matrix, so that the formation and the re-formation of the protective [70]

. A continuous silica layer was formed when the oxidation time was

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chromia layer was promoted

increased up to 1000 h (Fig. 11). Once the continuous layer of silica was formed between the oxide

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scale and substrate, the oxide scale would further impede the diffusion of iron cations from substrate to oxide scale as the lower diffusion rate of Fe in SiO2 than Cr2O3 [71-74]. When the diffusion rate of

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iron was impeded, the amount of iron cations in the scale would decrease. Consequently, the high oxidation rate of steel would be slowed down [69, 75, 76]. At the same time, the pore band was formed at the outer/inner layer interface as the outward diffusion flux of Fe cations could not meet the

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requirement of the growth of magnetite, as shown in Stage D. Moreover, the silica blocked the vacancy flow left by Fe outward diffusion to form micro-cavities and greatly enhanced the adhesive force of

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the chromia layer to the steel matrix [77]. In conclusion, the oxide scale microstructure of two F/M steels was characterized and the SIMP

steel showed better corrosion resistance than P91 steel as expected. Nevertheless, the mechanism is still not fully understood and further investigation is needed. In order to conclude on the application of the F/M steels in the SCWR, corrosion tests in flowing SCW with different oxygen concentrations for longer exposure time would be needed.

5.

Conclusions The oxidation behaviors of F/M SIMP and commercial P91 steels have been investigated in

flowing deaerated SCW at 700 °C under 25 MPa up to 1000 h. The oxide scales formed on the two steels were characterized and the reason that SIMP steel exhibited better SCW oxidation resistance, compared with P91 steel, was analyzed. The following main conclusions are as follows: (1) The oxide scale formed on both SIMP and P91 steels exhibited double layer structures consisting of an outer Fe3O4 layer and an inner Fe-Cr spinel layer, and there was an IOZ was not often observed

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due to the deaerated conditions. The thickness of the oxide layer of SIMP steel was smaller than that of P91 steel.

(2) Porosity caused by the outward migration of iron was observed in the inner layer, and primarily located at the spinel layer on SIMP steel and the inner layer-matrix interface on P91 steel. A porous

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band, with size of 10-15 µm, was formed at the oxide/metal interface of SIMP steel after exposure to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h.

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(3) The better oxidation resistance of SIMP steel than P91 steel was largely due to the higher Cr and Si content of the inner Fe-Cr spinel layer on SIMP steel and the 400 nm thick silica layer located at

steel in comparison with P91 steel. 1.

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the oxide/metal interface, which can be attributed to the higher chromium and silicon contents in SIMP

An additional reason that the SIMP steel showed better oxidation resistance was that the prior

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oxidized carbides, located at the diffusion layer, formed a continuous Cr/Si oxide mesh, impeding

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the outward migration of iron to form magnetite in comparison with P91 steel. Acknowledgments

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This work was financially supported by the project 2018NMSAKF03 of CAS Key Laboratory of

Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, the National Key Research and Development Program (No. 2016YFF0202200) and the sub-project (Nos. XDA03010301, XDA03010302) of Advanced Fission Energy Program-ADS Transmutation System, the Chinese Academy of Sciences Strategic Priority Research Program (No. XDA03010000).

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255-275.

Figure list

Fig.1. Weight gain of SIMP and P91 steels exposed to flowing deaerated SCW at 700 oC under 25

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Fig.2. XRD patterns of SIMP and P91 steels exposed to flowing deaerated SCW at 700 °C under 25

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Fig. 3. SEM images morphologies of the surface oxide scales formed on SIMP and P91steels exposed to flowing deaerated SCW at 700 °C under 25 MPa for different time (a, c, e, g, i) SIMP steel and (b, d, f, h, j) P91 for 200, 400, 600, 800 and 1000 h, respectively

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Fig. 4. Morphologies of surface oxide scales formed on SIMP steel exposed to flowing deaerated SCW at 700 °C under 25 MPa for 400 h

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Fig. 5. Morphologies of the exfoliated region surface oxide formed on SIMP steel exposed to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h

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Fig. 6. SEM images (a, b) and EDS analysis (c, d) of cross-section morphologies of oxide scales formed on SIMP and P91steels exposed to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h

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Fig. 7. SEM-BSE images and EDS elemental analysis of cross-sectional morphologies of oxide scales formed on SIMP steel exposed to flowing deaerated SCW at 700 °C under 25 MPa for 200 h

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Fig. 8. SEM-BSE images of cross section of oxide scale formed on SIMP and P91steels exposed to flowing deaerated SCW at 700 °C under 25 MPa for different time (a, c, e, g, i) SIMP steel and (b, d, f, h, j) P91 for 200, 400, 600, 800 and 1000 h, respectively

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Fig. 9. Oxide scale thickness and ratio of outer to inner layer thickness of SIMP and P91 steels exposed to flowing deaerated SCW at 700 °C under 25 MPa for different time

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Fig. 10. EPMA images of Fe-Cr spinel layers of SIMP and P91 steels exposed to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h

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Fig. 11. STEM images and elemental line scanning of interface between the Fe-Cr spinel layers and matrix of SIMP steel exposed to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h

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Fig. 12 Schematic representation of mechanism of oxide scale formation on SIMP steel during exposure to flowing deaerated SCW at 700 °C under 25 MPa

Table lists Table 1 Chemical compositions of experimental steels (wt.%) C

Si

Cr

Mn

W

Ta

V

Nb

Ni

Mo

P (ppm)

S (ppm)

P91

0.10

0.26

8.5

0.46

－

－

0.20

0.04

0.17

0.92

80

70

SIMP

0.22

1.22

10.24

0.52

1.45

0.12

0.18

0.01

－

－

30

40

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Steels

Table 2 Elemental composition at selected locations of SIMP steel exposed to flowing deaerated SCW at 700 °C under 25 MPa for 1000 h (wt.%) Mn Location Fe O Cr Si 1

74.36

25.64

--

--

--

2

72.20

26.40

1.41

--

--

3 4

75.02 53.65

21.76 18.47

2.02 23.51

-1.07

1.19

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3.30

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Table 3 Elemental contents of inner Fe-Cr spinel of SIMP and P91 steels exposed to flowing deaerated SCW at 700 °C under 25 MPa for different exposure times (wt.%) Steels Time (h) O Si Cr Fe 200 16.24 0.41 17.28 66.07 400 17.45 0.52 16.22 65.81 P91 600 16.36 0.39 16.84 66.41 800 16.93 0.54 17.06 65.47 1000 19.10 0.51 16.13 64.26 200 18.09 3.50 20.14 58.27 400 16.93 3.74 20.55 58.78 SIMP 600 19.64 3.54 20.24 56.58 800 16.86 3.35 20.52 59.28 1000 18.82 3.65 20.65 56.88