Corrosion and Stress Corrosion Cracking of Austenitic Alloys in Supercritical Water

Transcription

1 Corrosion and Stress Corrosion Cracking of Austenitic Alloys in Supercritical Water J. McKINLEY 1*, S. TEYSSEYRE 1, and G. S. WAS 1, D. B. MITTON 2, H. KIM 2, J-K KIM 2 and R. M. LATANISION 2, 1 Department of Nuclear Engineering and Radiological Sciences, College of Engineering, University of Michigan, 2355 Bonisteel Boulevard, Ann Arbor, MI , USA 2 Dept. Materials Science and Engineering, Massachusetts Institute of Technology 77 Massachusetts Avenue - Room 8-204, Cambridge, MA Abstract- The purpose of this study is to determine the corrosion and stress corrosion cracking (SCC) behavior of austenitic alloys for potential use as structural materials in the supercritical water reactor concept. SCC of 304L and 316L stainless steels was investigated using constant extension rate tensile experiments in either deaerated or non-deaerated supercritical water at 500 or 550 C. Corrosion experiments on heated tubes were conducted on 316L and Inconel 625 in high purity, nondeaerated water between 300 and 500 C. Results reveal that 304L is susceptible to intergranular SCC in 550 C nondeaerated water and, to a lesser extent, in 500 C deaerated water. The 316L sample failed by ductile rupture after 35% strain in 500 C deaerated water. Oxides formed on the SCC samples varied in thickness from ~5 microns in non-deaerated water to microns in deaerated water. The oxides were predominantly iron oxides with a significant amount of chromium oxides. The probable form of the oxides are FeO and Fe(OH) 3 and Cr 2 O 3, Cr(OH) 3, or CrOOH. In corrosion tests, the weight gain of 316L increases by a factor of three between 300 C and 500 C, with an oxide thickness of 5 microns after 8 days at 500 C. The oxide development in 625 is less than that in 316L. I. Introduction One of the most promising advanced reactor concepts for Generation IV nuclear reactors is the Supercritical Water Reactor (SCWR). Operating above the thermodynamic critical point of water (374ºC, 22.1 MPa), the SCWR offers many advantages compared to state-of-the-art LWRs including the use of a single phase coolant with high enthalpy, the elimination of components such as steam generators and steam separators and dryers, a low coolant mass inventory resulting in smaller components, and a much higher efficiency (~44% vs. 33% in current LWRs). Overall, the design provides a simplified, reduced volume system with high thermal efficiency. 1) Since supercritical water has never been used in nuclear power applications, there are numerous potential materials challenges that must be considered. Water in this phase exhibits properties significantly different from those of liquid water below the critical point. It acts like a dense gas and its density can vary with temperature and pressure from less than 0.1 g/cc to values similar to that of water below the critical point. This allows one to tune the properties of SCW, such as the ion product, heat capacity, and dielectric constant, to fit the application of interest. 2) The corrosion behavior of SCW over this range of densities varies widely depending upon the values of the properties present. 2-9) Supercritical water used as reactor coolant will fall in the lower end of the density scale at around 0.2 g/cc. At this density, water is a nonpolar solvent and can dissolve gases like oxygen to complete miscibility. Depending upon what species are present and how much oxygen is present in the solution, SCW in this state can become a very aggressive oxidizing environment. 2,4) This is a cause for concern about the general corrosion and stress corrosion cracking (SCC) of the structural materials and fuel elements of the reactors. The purpose of the current study is to investigate the corrosion and stress corrosion cracking of a variety of structural materials for use in SCW-cooled nuclear reactors. This paper focuses on two investigations. The first is a set of constant extension rate tensile experiments in SCW used to simulate stress corrosion cracking in 304L and 316L stainless steel. The stress corrosion cracking results were complemented by an investigation of the structure and composition of the oxide formed during the experiments. The second is a set of corrosion experiments in SCW performed over a range of temperatures on 316 stainless steel and Inconel 625. II. Experimental procedure 1. Materials Two commercial purity stainless steel alloys, 304L and 316L and one nickel base alloy 625 were used for this study. Both 304L and 316L are used for core internal components in light water reactors and the extensive database on the behavior of these alloys in BWR and PWR environments provides a reference condition against which the data in SCW can be compared. Inconel 625 has been used in reactor core and control rod components in LWRs as well as during the fabrication of bench-scale and pilot plant reactors used for waste destruction by SCW oxidation ) * Phone: (734) Fax: (734) jmckinle@engin.umich.edu 1

2 Table 1. Chemical compositions of the test alloys (wt%) used in the (a) corrosion experiments and (b) SCC experiments. (a) Elemental Analysis of Alloys Used in SCC Experiments Alloys C Mn Fe S Si Ni Cr Mo Cu N Co Nb P 316L Bal NM L Bal NM NM NM 0.02 (b) Elemental Analysis of Alloys Used in Corrosion Experiments Alloy C Mn Fe S Si Ni Cr Mo Cu N Co Nb P 316L Bal NM NM NM NM Bal NM NM: Not measured Stainless steel alloys 316L and 304L were studied in the SCC experiments, and the stainless steel alloy 316L and Inconel alloy 625 were studied in the general corrosion experiments. The composition of each of these alloys is shown in Table 1. The alloys for the SCC experiments were first heat treated and then machined into SCC bars 38 mm long with threaded ends and a gage dimension of 2mm x 1.5mm x 21mm. The 304L alloy was used in the as-received condition, which contained a grain size of 40 µm. The 316L alloy was solution annealed at 1100 C for 20 minutes and water quenched to remove any carbide present. The grain size after this heat treatment was 44µm. SCC bars were machined from the heat-treated material by electric discharge machining. Both alloys were mechanically polished using standard metallographic techniques and then electropolished to obtain a mirror finish. 13) Alloys for corrosion experiments were obtained in tube form (approximately 3.2 mm OD and 1.1 mm ID) and used in the as-received condition. Prior to an experiment, the tube inside diameter was cleaned with acetone and subsequently rinsed with 15 MΩ deionized water. 2. Stress Corrosion Cracking Experiments Constant extension rate tensile experiments (CERT) were performed through use of an enclosed, flowing water loop schematically shown in Fig. 1. Dissolved oxygen in deionized water was controlled by Ar deaeration. Experiments were conducted at either O 2 <10 ppb (deaerated condition) or at O 2 ~8 ppm (non deaerated condition). The flow rate used was 10 ml/min. The conditions for each of the experiments are given in Table 2. The water is pressurized by an HPLC metering pump and heated in a convection preheater before it flows into the test vessel. The water under supercritical conditions flows back out of the autoclave and is cooled by a tube-in-tube heat exchanger before the pressure is reduced by the backpressure regulator (BPR). The water passes though an ion exchanger before flowing back into the main column. The conductivity and the oxygen content of the circulating water are measured at room temperature and atmospheric pressure in both the inlet and outlet lines. The temperature is measured in the vessel and the pressure is determined at room temperature in the inlet and outlet lines. The nominal oxygen content in a supercritical watercooled reactor will likely be very low (<10 ppb). However, the generation of oxygen and hydrogen gas by radiolysis and the high solubility of these gases in the supercritical water could support a higher concentration of gases during operation. Because of this, it was considered important to perform experiments in both deaerated and non-deaerated water. The sample is strained at a nominal rate of 3x10-7 s -1 using a stepping motor. The load train is configured with a pressure-balanced system to ensure that the sample is not under stress from the water pressure. This is necessary in order to avoid sample yielding due to the high internal pressure. After completion of the experiment, the samples were examined using a Philips XL-30 scanning electron microscope (SEM). The fracture surface and the faces of the sample gages were analyzed. 3. Analysis of Oxide Formed During the Stress Corrosion Cracking Experiments The oxides formed on the samples were examined using a variety of techniques. The fracture surfaces and the faces of the sample gage section were examined using a Philips XL-30 field emission gun scanning electron microscope (SEM). The composition of the oxide on the surfaces of the samples was analyzed using energy dispersive spectroscopy (EDS) in conjunction with the SEM. These studies were carried out at an accelerating voltage of 15 kv and a working distance of 10mm. Monte Carlo simulations were used to determine the depth of penetration of the electron beam into the oxide. These calculations were compared with the oxide thickness measured on cross-sections taken of the oxidized SCC bars. Even the thinnest oxide layer was thick enough to prevent the electron beam from penetrating to the metal substrate. The oxide was analyzed using X-ray photoelectron spectroscopy (XPS) on a Philips Phi 5400 XPS system to 2

3 determine the bonding state of the elements present in the oxide. The scans were performed using an Al Kα X-ray source (E hυ = ev) operating at a power of 300 W and using a hemispherical analyzer. The sample was cleaned by sputtering with an argon ion beam of energy 3 kev for one minute prior to analysis. This assured that much of the contamination present on the surface in the form of water and carbon compounds was removed. The spectra were deconvoluted by obtaining reference data for the various compounds expected to be present and anchoring the peaks used in the fit at or very near those reference points. The peak height and width were then used as fitting parameters. Figure 1: Schematic of the circulating SCW CERT system Table 2: Conditions for the flowing SCW CERT tests. Alloy 304L 316L Environment SCW SCW Argon SCW Temperature ( C) Pressure (MPa) O 2 (ppb) 6-8 ppm < < 10 Inlet conductivity (µs/cm) <0.5 < < 0.1 Flow rate (ml/min) Strain rate (s -1 ) variable 1 3X10-7 3X10-7 3X10-7 Test duration (hr) Strain to failure (%) Stress at failure (MPa) Yield strength (MPa) Failure mode Ductile rupture initiated by intergranular fracture Intergranular cracks initiated Ductile Ductile 1 The strain rate was 1 x 10-6 s -1 for the first 4.25% strain and 5 x 10-7 s -1 for the balance of the experiment. 2 Test was stopped prior to failure 3

4 Thermocouples Flow Cool Down Heat Exchangers Computer Controller Instrumented Tube Cooling Water Autoclave Heater Back Pressure Regulator Effluent Preheater Preheater Flow Measurement DI Water Pump Flow Pump Flow DI Water Lexan Shield Figure 2: Schematic drawing of the flowing supercritical test system for studying corrosion at MIT 4. Corrosion Experiment Figure 2 presents a schematic representation of the current supercritical corrosion system configuration at MIT, including a traditional alloy 625 autoclave (right of figure) and an instrumented tube (left of figure) system. The temperature of the feed stream is increased to supercritical temperature by a preheater. Pressure is provided by a back pressure regulator (BPR) and an HPLC pump. The temperature and pressure are controlled by a computer and lexan shields are provided for safety. When the autoclave system is in use, standard mass loss or u-bend samples are mounted on a rack and inserted into the autoclave. Subsequent to an experiment samples are removed and assessed both metallographically and analytically. In the case of the instrumented tube experiments, a tube of the alloy to be tested is used as the autoclave, and micro-thermocouples are attached externally along the length of the vessel. Water at an elevated temperature and pressure is pumped into one end and permitted to cool as it traverses the tube. In general, once steady state is achieved, temperature fluctuations are minor (+/- 2 C). Water-cooling by external copper tubes is incorporated between the final two thermocouples in order to achieve the significant temperature drop needed to assess both sub- and supercritical temperatures simultaneously. On completing a test, the tube is sectioned, mounted and examined. This experimental configuration is used to correlate temperature to corrosion rate and mode. III. Results 1. Stress Corrosion Cracking Experiments Constant extension rate tensile experiments were performed on commercial purity 304L and 316L in supercritical water and in argon at 500 and 550 C. The 304L sample pulled in deaerated SCW was strained until 25% strain, and the other samples were pulled to failure. Experimental conditions and results are summarized in Table 2. The stress-strain curves for the 304L samples are plotted in Fig. 3a and those for the 316L samples in Fig. 3b. In all experiments in supercritical water, the stress increased initially without movement of the pull rod because of the static friction between the bal seals and the pull rod. Once the applied stress exceeded the static friction stress, the pull rod started to move and the stress read by the load cell dropped to the value of the dynamic friction stress. The sample began to strain after the slack in the load train was eliminated. The stress recorded under supercritical conditions was noisy and the variation reached as much as ±20MPa. Up to 22% strain, the stress-strain behaviors of the 304L samples in non-deaerated and deaerated water were similar. The test in non-deaerated water proceeded normally and resulted in failure at 32% strain. The sample in deaerated water was terminated after 25% due to a problem with the pump. The stress discontinuity at 22% strain was due to an interruption in the test at that strain level. Both 316L samples proceeded to failure without major 4

5 500ºC, 25.5 MPa Deaerated water Conductivity: <0.1 µs/cm Strain rate: 3 x 10-7 s L 500ºC, Argon Strain rate: 3 x 10-7 s L 550ºC, 25.5 MPa Non-deaerated water Conductivity: <0.5 µs/cm: Strain rate: 5 x 10-7 s ºC, 25.5 MPa Deaerated water Conductivity: <0.1 µs/cm Strain rate: 3 x 10-7 s -1 a) b) Figure 3: Stress-strain behavior of a) 304 in SCW in non-deaerated and deaerated conditions, and b) 316L in deaerated SCW and in argon. interruptions. The 316L sample strained in deaerated supercritical water exhibits a behavior similar to that of the 316L sample strained in argon. The 316L sample strained in argon showed a slightly higher maximum stress and maximum strain. The 304L SS sample strained in non-deaerated SCW displayed evidence of intergranular stress corrosion cracking (IGSCC). In this sample, necking of the cross section was minimal as shown in the micrograph in Fig. 4a. A large number of secondary cracks were visible on the side of the sample such as the one shown in Fig. 4b. The density of the cracks was around 20 cracks/mm 2. Cracks also occurred well away from the fracture region. At high magnification, the secondary cracks clearly appeared intergranular in nature. Figure 4c shows the fracture surface at the corner of the sample. Intergranular fracture occurred in multiple locations near the edge and corners of the fracture surface. The balance of the fracture surface was typical of ductile rupture, indicating that failure occurred by this mode but was initiated by intergranular fracture. It should be noted that this same heat of 304L has previously been tested under PWR conditions (320 C, deaerated water containing 2 ppm Li and 1000 ppm B, <5 ppb O 2, 35cc/kg H 2 and 20.5 µs/cm), BWR normal water conditions (2000 ppb O 2 and 0.2 µs/cm, +140 mv SHE ), and hydrogen water chemistry conditions (<5 ppb O 2, 560 ppb H 2 and 0.1 µs/cm). Intergranular cracking was not observed on either the fracture surface or the side surface of the samples tested under these conditions. The 304L SS sample strained in deaerated water reached 25% strain before the test was stopped. Subsequent examination revealed that the sample did not neck, similar to the behavior observed in non-deaerated SCW. Also, there were a significant number of cracks on the sample surface (see Fig. 5a). Inspection of the cracks at high magnification (see Fig. 5b) revealed that they were intergranular cracks. The crack density was about 7 cracks/mm 2, or approximately one-third of the magnitude of cracking observed on the 304L sample tested in aerated water. The lower oxygen content used for the deaerated sample resulted in a less oxidizing environment and may be the cause of the lower crack density. However, the observation of IGSCC in 304L in both the non-deaerated and deaerated conditions is significant. The deformation behavior of the 316L alloy was similar in argon and deaerated SCW. In both experiments, there was significant necking of the samples. Both samples showed dimpled, ductile fracture on the fracture surface and no evidence of intergranular cracking on the sample surfaces, as shown in Fig. 6. In deaerated SCW, alloy 316L did not display evidence of IGSCC. 304L, however, displayed intergranular cracking in both non-deaerated and deaerated supercritical water. Since the initiation of cracks in 304L appeared to increase with increasing oxygen content of the SCW, the question of the possibility of observing IGSCC in 316L by increasing the SCW oxygen content is still open. 5

6 b) a) c) Figure 4: 304 sample after failure in non deaerated SCW; a) and b) the surface of the sample exhibits a high density of secondary intergranular cracks, and c) intergranular fracture on the fracture surface near the corner of the sample initiated failure in the sample before significant necking occurred. b) a) Figure 5: The 304L sample strained to 25% in deaerated SCW: a) the surface of the sample showing a high density of cracks, and b) a crack at high magnification showing the intergranular nature of the crack. 6

7 a) b) Figure 6: 316 sample after failure in deaerated SCW: a) the surface of the sample showing necking near the rupture point, and b) fracture surface showing a ductile fracture. 2. Results of the Analysis of the Oxide Formed During the Stress Corrosion Cracking Experiments The oxides formed on the 304L and 316L stainless steel samples were analyzed in further detail using a combination of sample cross-sections, energy dispersive spectroscopy (EDS), and X-ray photoelectron spectroscopy (XPS). These techniques provided information about the oxide thickness, the composition, and the binding state of the compounds present in the oxide. (1) Sample Cross-Sections Table 3 presents oxide thicknesses calculated from scanning electron microscope images like those shown in Fig. 7 along with the conditions under which the experiments were performed. These oxide thicknesses show the expected behavior in most cases. The 304L exposed to non-deaerated SCW formed a much thicker oxide layer than did the 304L sample exposed to deaerated SCW, which coincides with the expected presence of a more oxidizing environment in the former. The 316L sample exposed to deaerated SCW had a slightly thinner oxide than that on the 304L sample exposed to deaerated SCW. The thickness of the oxide layer formed on the 316L sample exposed to argon at 500 C is higher than was first expected. This sample was exposed to deaerated SCW for five days before being placed in an argon atmosphere while a mechanical problem with the SCW system was being fixed. This experimental step would have explained the presence of an oxide layer of a comparable thickness to that of the 316L sample exposed to deaerated SCW if such had been observed. However, the presence of an oxide layer formed in argon which was thicker than that of the oxide layer formed in deaerated SCW indicated that the flowing argon was not purified from oxygen enough to limit oxidation. (2) Energy Dispersive Spectroscopy of the Oxides The results of the EDS measurements taken on each sample are presented in Table 4. The values in Table 4 are averages of the data taken by the EDS for each sample. These results show that the part of the oxide layer that is being measured is composed primarily of iron oxides and contains proportionately less chromium and nickel than does the metal matrix. The 304L sample exposed to aerated SCW is an exception in that it shows a much higher concentration of chromium in the oxide film. The fact that the proportion of iron to chromium in the oxide was of this magnitude Table 3: Results of oxide thickness measurements compiled with the experimental conditions. Conditions Alloy Oxide Thickness (µm) Temperature ( C) Pressure (MPa) Environment Exp. Time (hrs) 304L Non-deaerated SCW L Deaerated SCW L Argon L Deaerated SCW Under atmospheric pressure. 7

8 Figure 7a: SEM backscattered electron image of a crosssection of 304L exposed to non-deaerated SCW for 160 hrs at 550 C and 25.5 MPa. Figure 7b: SEM backscattered electron image of a crosssection of 304L exposed to deaerated SCW for 230 hrs at 500 C and 25.5 MPa. 5000x 316_base_1 5000x Figure 7c: SEM backscattered electron image of a crosssection of 316L exposed to argon at atmospheric pressure for 320 hrs at 500 C after 5 days in deaerated SCW. Figure 7d: SEM backscattered electron image of a crosssection of 316L exposed to deaerated SCW for 305 hrs at 500 C and 25.5 MPa. Table 4: Results of EDS measurements on the oxide formed on the SCC sample bars. Composition Measured by EDS 1 (at%) Sample Experimental Conditions Fe Cr Ni O 304L 550 C 25.5 MPa non-deaerated SCW L 500 C 25.5 MPa deaerated SCW L 500 C 25.5 MPa deaerated SCW Compositions normalized to 100% for the elements of interest. Carbon contamination is not included. indicates that the environment stabilizes iron oxides more than chromium or nickel oxides. To confirm that the EDS scans were measuring only the oxide and not the substrate, Monte Carlo calculations were made for the electron beam oxide interaction. The calculations for a 15kV electron beam determined that the X-rays measured by the EDS could penetrate an oxide thickness of no more than one micron. All of the samples analyzed exhibited oxides thick enough to prevent the beam from including the metal substrate in the EDS measurement. X-ray photoelectron spectroscopy (XPS) was used to determine the bonding state of the oxide. The shallow depth 8

9 of penetration of the x-ray beam essentially guarantees that only the very top layer of the oxide is surveyed for data on composition and bonding. (3) X-ray Photoelectron Spectroscopy of the Oxides XPS survey scans for the 304L sample exposed to 550 C non-deaerated SCW and for the 316L sample exposed to 500 C deaerated SCW are shown in Fig. 8. Both samples exhibited mostly iron, chromium, and oxygen in the survey scans of the oxide. The 316L also showed a small peak in the nickel region. This was nearly negligible for the 316L sample and did not show up at all for the 304L sample. The carbon peak is not considered in analysis or calculations because the majority of it is due to surface contamination from carbon dioxide, hydrocarbons, and other carboncontaining contaminants. Results of focused region scans for the two oxidized samples are presented in Table 5. Each sample was scanned in detail in the oxygen 1s, the iron 2p, the chromium 2p, and the nickel 2p regions. Table 5 shows the results of the peak fitting done for each of the oxidation states found present in the oxide. The peak locations were obtained from a review of various references, and thus correspond well to tabulated values. The focused region scans for the oxidized 316L sample are presented in Fig. 9. The scans for the oxidized 304L sample were very similar and are not presented. These results provide a number of observations about the oxidation behavior of these alloys in SCW. First, none of the focused scans reveal base metal to be present. All of the nickel, iron, and chromium present is in an oxidized or hydroxylated state. Second, there is a negligible amount of nickel oxide present. The alloys contain about 8-10% Ni, but the rate at which the nickel oxidizes appears to be small Figure 8a: XPS survey scan of the oxide on the 304L sample exposed to non-deaerated SCW for 160 hours at 550 C and 25.5 MPa. Figure 8b: XPS survey scan of the oxide on the 316L sample exposed to deaerated SCW for 305 hours at 500 C and 25.5 MPa. Table 5: XPS results for the oxides formed during two SCC experiments Alloy Oxidation States Present for each Component Experimental Fe 2+ Fe 3+ Cr 3+ Ni Conditions 2+ Peak (ev) Area Peak (ev) Area Peak (ev) Area Peak (ev) Area 304L 550 C 25.5 MPa nondeaerated SCW N/A N/A 316L 500 C 25.5 MPa deaerated SCW

10 Figure 9a: XPS results for oxidized 316L in the iron 2p region. The curve fit includes Fe 2+, Fe 3+, and the corresponding satellite peak. The iron metal peak at ev is not present. Figure 9b: XPS results for oxidized 316L in the chromium 2p region. The chromium is present in the Cr 3+ state. The chromium metal peak at ev is not present Figure 9c: XPS results for oxidized 316L in the nickel 2p region. The nickel is present in the Ni 2+ state. The nickel metal peak at is not present. Figure 9d: XPS results for oxidized 316L in the oxygen 1s region. The oxygen is present in OH - and O 2- form. The Water and other peak corresponds to adsorbed water and gas contamination on the surface. 10

11 enough compared to the rate of iron and chromium oxidation to make it a negligible contributor to the oxide layer. The most important result is the bonding state of the oxides. Both iron and chromium can exhibit a variety of oxidation states. In these samples, iron appeared to take on the Fe 2+ and Fe 3+ states while chromium took on only the Cr 3+ state. The Fe 2+ peak at ev probably corresponds to FeO and the Fe 3+ peak at ev probably corresponds to either FeOOH or Fe(OH) 3. 14,15,16) These peaks could also represent the presence of Fe 2 O 3 and Fe 3 O 4. However, the satellite peak for Fe 2 O 3 would be present at 720 ev if this compound were present, and if Fe 3 O 4 were present in any significant amounts the satellite peak indicated above would be smaller since magnetite does not exhibit a satellite peak. 16) The values of both the 2p 3/2 and the 2p 1/2 peak correspond to various reference values for FeO and FeOOH or Fe(OH) 3. 14,15,16) The area ratio between the iron present as an oxide and that present as a hydroxide is approximately equal in magnitude. The Cr 3+ peak at ev could correspond to either Cr 2 O 3, Cr(OH) 3, or CrOOH. 17) The 2p 1/2 peak present in these data sets is higher than the values normally seen for Cr 2 O 3 of around ev. 18) CrOOH peak values are generally higher than that of Cr 2 O 3 and thus would correspond better with the 2p 1/2 values of ev for the 316L oxide and ev for the 304L oxide. Thus the chromium present in the oxide layer is probably in the form of a hydroxide. The oxygen 1s peak can provide a bit more data about the bonding state of the metal oxides. For both of the samples the oxygen peak was fit to a hydroxide curve, an O 2- curve, and a curve to pick up any adsorbed water or oxygencontaining gases. This third peak eliminates the possibility of doing quantitative analysis on the oxygen peak, but the presence of both oxygen states in the layer corresponds well with the information provided by the metal peaks. Having the metals bound to oxygen in both the O 2- state and the OH - state makes physical sense when considering high temperature aqueous systems. The oxygen that is present in the system as a dissolved gas will react with the metal to form the O 2- state bonds. The hydroxide ions would come from the dissociation of water molecules. This dissociation could produce some O 2- ions, but it is much more energetically favorable for the reaction to produce hydroxide ions which could then react with the metal. 3. Corrosion Experiments A schematic of the experimental configuration and micrographs of the 316L stainless steel tube cross section at several temperatures are presented in Fig. 10. The sample was exposed for approximately 192 hours to non-deaerated 15 MΩ water at a pressure of 24 MPa. The highest temperature achieved was in excess of 395 C. The preliminary analysis of the tube cross-section suggests substantial oxide development at the highest temperature, with decreasing oxide development at lower temperatures. 200ºC 270ºC 338ºC 348ºC >395ºC Figure 10: Experimental configuration and micrographs of the 316L stainless steel tube cross section at several temperatures after exposure to non-deaerated 15 MΩ water for approximately 192 hours. 11

12 The inner diameter of the tube revealed oxide development of approximately 5 µm at the highest temperature assessed. That the tube reflects a decrease in the extent of oxidation as a function of decreasing temperature was in agreement with tests on mass loss coupons. After extended exposure (10-12 days), type 316 stainless steel exhibited an average mass gain (standardized to area and exposure time) on the order of mg/cm 2 /day at a temperature of 300 C. This value increased to mg/cm 2 /day at 500 C. The elemental analysis of the 316 tube cross section presented in Fig. 11 indicated the presence of an oxide layer; however, the oxide has not yet been precisely identified. The sample does not reveal the preferential elemental dissolution previously seen in SCWO systems. The experimental configuration and micrographs of the alloy 625 tube cross section at several temperatures are presented in Fig. 12. This sample was exposed for approximately 210 hours to non-deaerated 15 MΩ water at a pressure of 3500 psig. The temperature range covered was approximately 425 to 300ºC. O Fe Cr Ni Mo O Figure 11: Elemental analysis of the 316 stainless steel tube cross section exposed to non-deaerated 15 MΩ water for approximately 192 hours at a temperature greater than 395 C. 300ºC 343ºC 357ºC 365ºC 419ºC Figure 12: Experimental configuration and micrographs of the alloy 625 tube cross section at several temperatures after exposure to non-deaerated 15 MΩ water for approximately 210 hours 12

13 Ni Cr O Figure 13: Elemental analysis of the alloy 625 tube cross section exposed to non-deaerated 15 MΩ water for approximately 210 hours at a temperature greater than 419 C. Analysis of the 625 tube revealed substantially less oxide development in the high temperature region than was seen for the stainless steel tube sample. Again, this was in agreement with previous work. 5) Surface analysis by EDX mapping is shown in Fig. 13 and indicated that an oxide was present. While the oxide has not yet been identified, the indication was that the chromium level was relatively high. Unlike previous work that was carried out in more aggressive environments, 6-12) neither preferential dissolution nor crack development have been observed to date for the corrosion samples. IV. Conclusion Stainless steel alloy 304L was found to be susceptible to intergranular stress corrosion cracking in both nondeaerated, high conductivity supercritical water and deaerated, low conductivity supercritical water at 550 C and 500 C, respectively. The extent of the IGSCC was the most severe for the higher temperature, non-deaerated condition. The particular 304L alloy heat used in these experiments is resistant to IGSCC in simulated BWR normal water chemistry at 288 C and in simulated primary water at 320 C. The 316L stainless steel did not exhibit any evidence of IGSCC in 500 C deaerated supercritical water or in argon at the same temperature. Both 316L samples achieved approximately 35% total elongation prior to ductile, dimpled rupture. Oxide cross-section analyses revealed that the 304L sample in a non-deaerated SCW environment formed a thicker oxide layer than samples exposed to deaerated environments. Cross-sections also revealed that the thickness of the oxide formed on 304L after 160 hours in 550 C non-deaerated SCW was equivalent to that of the oxide observed on 316L after exposure to non-deaerated SCW for 192 hours at 400 C. Analysis of the oxide chemistry revealed that the oxide layers were primarily made up of iron and chromium oxides and hydroxides. The iron oxides and hydroxides made up the majority of the composition and were present in approximately a one to one ratio to one another. These results for the very top layer of the oxide indicate that the iron oxides and hydroxides form at a faster rate than the chromium or nickel oxides and hydroxides. This would imply that the iron oxides form first but do not protect the metal surface well enough to significantly affect the corrosion rate. The chromium oxides could then begin to form and eventually create a protective oxide coating to slow the overall oxide growth. To confirm this hypothesis further XPS and EDS testing of the oxide layer as a function of depth would be required. Corrosion experiments of 316L stainless steel in high purity, non-deaerated water revealed a strong dependence of oxide formation on the temperature of the environment. Experimental results showed that 316L forms an oxide at 500 C at a rate three times greater than at 300 C. The thickness of this oxide was about 5 microns after 8 days at 500 C. Inconel alloy 625 was investigated over a temperature range of C and showed less oxide development than the 316L stainless steel sample V. Acknowledgements The authors wish to thank Dr. Yongsun Yi for his role in the design and fabrication of the SCWR test loop in the High Temperature Corrosion Laboratory at the University of Michigan. They also wish to thank Brent Capell for his assistance with obtaining and interpreting the XPS data. This work was supported by a contract with Bechtal Power Corporation and a prime contract from the U. S. Department of Energy under the Nuclear Energy Research Initiative program, project #

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