HYDROGEN EMBRITTLEMENT EVALUATION METHOD FOR Cr-Mo-V STEEL

Abstract

To appropriately evaluate hydrogen embrittlement of members including Cr-Mo-V steel.SOLUTION: An evaluation method includes: a step of obtaining information of a structure of a material and information of physical properties, start-up in using high-temperature high-pressure hydrogen, steady operation, operation time at shutdown, a process temperature according to operation lapsed time, and a process pressure and information of hydrogen concentration under such an environment; a step of simulating, on the basis of each information, a temperature generated during operation by members, and phenomena of a stress distribution and a diffusible hydrogen distribution across a wall thickness; a practical property providing step of subjecting a sample material obtained by reflecting a member to be evaluated or having crack-like failure to manufacturing conditions executed when manufacturing the member to provide a material property of a practical material; classifying phenomena in use into start-up, steady operation, a temperature generated at shutdown, stress and a diffusible hydrogen content; and reflecting these conditions in the sample material to evaluate hydrogen-enhanced cracking development.SELECTED DRAWING: Figure 45

Description

[Technical field]

[0001] The present invention relates to a method for evaluating hydrogen embrittlement of Cr-Mo-V steel used in oil refining reactors and the like. [Background technology]

[0002] Hydrogen embrittlement causes crack propagation in materials used in oil refining reactors in hydrocracking and desulfurization processes. This crack propagation is classified into two types: Internal Hydrogen Assisted Cracking (IHAC), which occurs when diffusible hydrogen is present in the material, and Hydrogen Environment Assisted Cracking (HEAC), which occurs when the tip of the crack is exposed to H2 gas and hydrogen penetrates into the material.

[0003] It has been shown that 2.25Cr-1Mo steel reactors used in reactors are prone to internal hydrogen-assisted cracking during shutdown, and that these properties may depend on the hydrogen concentration, the degree of temper embrittlement, loading rate (dK / dt), temperature, and the geometry / microstructure of the steel. Furthermore, it has been suggested that the presence of gaseous hydrogen may degrade the steel through hydrogen environment assisted cracking (HEAC), including the possibility of interaction with IHAC (Non-Patent Document 1). ASME SA-336 Grade F22, Class 3 (forged material) and ASME SA-387 Grade 22, Class 2 (plate material) are commonly used as 2.25Cr-1Mo steel reactor materials. [Prior art documents] [Non-patent literature]

[0004] [Non-Patent Document 1] Richard P. Gangloff, Ph.D. and Ted L. Anderson, Ph.D., PE, The Effects of Hydrogen for Establishing a Minimum Pressurization Temperature (MPT) for Heavy Wall Steel Reactor Vessels, WHITE PAPER ON HYROGEN EFFECTS ON MPT, American Petroleum Institute, Published on February 21, 2018 Summary of the Invention [Problem to be solved by the invention]

[0005] Currently, 2.25Cr-1Mo-V steel is widely used in reactors. It is known that HEAC and IHAC phenomena occur in 2Cr-1Mo-V steel reactors, but evaluation of hydrogen embrittlement reflecting actual equipment cannot accurately predict the lifespan by evaluation only at shutdown, as is done with 2.25Cr-1Mo steel reactors. Therefore, it is desirable to accurately evaluate material properties to predict aging deterioration and damage and to evaluate fitness for service. 2.25Cr-1Mo-V steel reactor materials generally used are ASME SA-336 Grade F22V (forged material) and ASME SA-542 Type D, Class 4a (plate material).

[0006] The present invention has been made against the above background, and has an object to provide a method for evaluating hydrogen embrittlement that can appropriately evaluate the hydrogen embrittlement of Cr-Mo-V steel used in a hydrogen environment. [Means for solving the problem]

[0007] In one embodiment, the phenomena occurring in a Cr-Mo-V steel reactor during its operation, such as temperature, stress, and diffusible hydrogen distribution, are clarified through simulation, and the obtained hydrogen embrittlement factors are reproduced in a rising load test using test material that reflects the actual equipment material, making it possible to evaluate the fracture mechanics characteristics of the reactor against hydrogen-assisted cracking propagation. Effect of the Invention

[0008] According to the present disclosure, it is possible to properly evaluate hydrogen embrittlement of components containing Cr-Mo-V steel, and it becomes possible to set the minimum pressurization temperature for hydrogen-assisted cracking growth and to evaluate the life of the components. [Brief description of the drawings]

[0009] [Figure 1] 1 is a graph showing temperature and pressure versus start-up and shut-down time for a 2.25Cr-1Mo-V reactor according to an embodiment. [Diagram 2] FIG. 1 is a diagram showing the basic design structure of a 2.25Cr-1Mo-V reactor in an embodiment. [Diagram 3] 13 is a graph showing the time history response of hydrogen concentration on the inner surface, the boundary WOL side, and the 2.25Cr-1Mo-V side in response to startup and shutdown of a reactor in an embodiment. [Figure 4] This figure shows the basic dimensions, components and materials of the 6”ID quench nozzle to be analyzed. [Diagram 5] This figure shows the specifications and boundary conditions (unsteady heat conduction analysis / hydrogen diffusion simulation) of the analysis model of the 6”ID quench nozzle. [Figure 6] 1 is a graph showing the time history response of temperature and hydrogen accumulation (results of heat transfer analysis and hydrogen diffusion simulation) for the start-up and shutdown of a reactor. [Figure 7] 1 is a graph showing the time history response (structural analysis results) of circumferential and axial stress components for reactor startup and shutdown. [Figure 8]1 is a graph showing the tensile strength at a design temperature of 454°C as a function of the PWHT (705°C) holding time of 2.25Cr-1Mo-V steel. [Figure 9] 1 is a graph showing a comparison of temper embrittlement resistance properties of 2.25Cr-1Mo steel and 2.25Cr-1Mo-V steel (transition temperature of Charpy impact absorbed energy of 55 J before and after step cooling). [Figure 10] FIG. 1 shows the groove shape of the shell ring circumferential welded joint of an actual reactor (left) and the groove shape of the 2.25Cr-1Mo-V steel welded test plate (right). [Figure 11] 1 is a graph showing the Charpy impact properties of the weld metal of a 2.25Cr-1Mo-V steel welded test plate. [Figure 12] This is a schematic diagram of the progress of internal hydrogen assisted cracking (IHAC) in a reactor. [Figure 13] 1 is a graph showing the relationship between the time left in air at room temperature and the hydrogen content depending on the hydrogen exposure temperature. [Figure 14] 1 is a graph showing the relationship between the exposure time in air at 150° C. and the hydrogen content depending on the analytical sample collection position. [Figure 15] FIG. 1 is a diagram showing the relationship between the thermal history and the sampling position of a hydrogen analysis sample and the hydrogen content. [Figure 16] This is a schematic diagram of the progress of reactor hydrogen environment assisted cracking (HEAC). [Figure 17] 1 is a graph showing the results of a rising load test (high toughness weld) in 20 MPa H2 gas at room temperature with and without hydrogen charging to a 1T-C(T) test piece. [Figure 18] These are micrographs of SEM observations of fracture surfaces (high-toughness welds) of 1T-C(T) test pieces after a rising load test in 20 MPa H2 gas at room temperature with and without hydrogen charging. [Figure 19] 1 is a graph showing the results of a rising load test (low toughness weld) in room temperature 20 MPa H2 gas with and without hydrogen charging of a 1T-C(T) test piece. [Figure 20] These are micrographs of the fracture surface SEM observation results (low toughness weld) of a 1T-C(T) test piece after a rising load test in 20 MPa H2 gas at room temperature with and without hydrogen charging. [Figure 21] Schematic diagram of crack growth behavior (IHAC+HEAC) during startup with hydrogen charging (room temperature). [Figure 22] 1 is a graph showing the results of a rising load test (high toughness weld) in 150° C., 20 MPa H2 gas with and without hydrogen charging on a 1T-C(T) test piece. [Figure 23] 1T-C(T) is a photomicrograph showing the results of SEM observation of the fracture surface (high toughness weld) after a rising load test in 20 MPa H2 gas at 150°C with and without hydrogen charging of a test piece. [Figure 24] 1 is a graph showing the results of a rising load test (low toughness weld) in 20 MPa H2 gas at 150°C with and without hydrogen charging on a 1T-C(T) test piece. [Diagram 25] These are micrographs of SEM observations of fracture surfaces (low toughness welds) of 1T-C(T) test pieces after a rising load test in 20 MPa H2 gas at 150°C with and without hydrogen charging. [Figure 26] Schematic diagram of crack growth behavior (IHAC+HEAC) for start-up with hydrogen charging (150°C). [Figure 27] 1 is a graph showing a load-load line displacement curve (high toughness weld) in a rising + holding load test in 20 MPa H2 gas at room temperature using a hydrogen-charged 1T-C(T) test piece. [Figure 28] 1 is a graph showing the history of load line displacement (high toughness welded portion) against the pseudo cleavage crack propagation initiation load holding for 10 days. [Figure 29] 1 is a graph showing a load-load line displacement curve (low toughness weld) in a rising + holding load test in 20 MPa H2 gas at room temperature using a hydrogen-charged 1T-C(T) test piece. [Diagram 30]1 is a graph showing the history of load line displacement (low toughness welded portion) against the pseudo cleavage crack propagation initiation load holding for 10 days. [Diagram 31] 1 is a micrograph showing a crack growth measurement with respect to a pseudo-cleavage crack growth initiation load hold. [Diagram 32] 1 is a graph showing the correlation between the crack length increment and the load line displacement increment with respect to the pseudo-cleavage crack growth initiation load hold. [Diagram 33] 1 is a graph showing a comparison of pseudo-cleavage crack growth rate models for high and low toughness welds. [Diagram 34] 1 is a graph showing the results of a rising load test (high toughness weld) in air at 150° C. and at room temperature after hydrogen charging of a 1T-C(T) test piece. [Diagram 35] These are micrographs of SEM observations of fracture surfaces (high-toughness welds) of 1T-C(T) test pieces after hydrogen charging and rising load tests in air at 150°C and room temperature. [Diagram 36] 1 is a graph showing the results of rising load tests (low toughness welds) in air at 150° C. and at room temperature after hydrogen charging of a 1T-C(T) test piece. [Figure 37] These are micrographs of the fracture surface SEM observation results (low toughness weld) of a 1T-C(T) test piece after hydrogen charging and rising load testing in air at 150°C and room temperature. [Figure 38] Schematic diagram of crack propagation behavior during shutdown (IHAC) under hydrogen charging conditions. [Figure 39] 1 is a graph showing predictions of the K values ​​(KIH and KIC-H) that define the initiation of hydrogen-assisted crack propagation and rapid fracture. [Diagram 40] 1 is a graph showing hydrogen-assisted crack propagation evaluation results (time history response of stress intensity factors for reactor startup and shutdown). [Diagram 41] 1 is a graph showing the results of a determination as to whether or not a reactor can be continued to be operated in the event of hydrogen embrittlement accompanied by the progress of hydrogen-assisted cracking. [Diagram 42] 1 is a micrograph showing the predicted subcritical crack growth rate from the onset of pseudo-cleavage crack growth (KIH value) to cleavage fracture (KIC-H value). [Diagram 43] 1 is a graph showing prediction of sub-critical crack growth rate from the initiation of pseudo-cleavage crack growth (KIH value) to cleavage fracture (KIC-H value). [Diagram 44] 1 is a graph showing the results of a hydrogen-assisted crack growth analysis (crack growth curve and stress intensity factor distribution) for repeated reactor operation. [Diagram 45] 1 is a graph showing the reactor remaining life evaluation results based on hydrogen-assisted crack propagation analysis (trajectory of crack propagation evaluation points due to da / dt to FAD). DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0010] An embodiment of the present disclosure will be described below. 1. Clarification of the phenomenon for the operation of a real reactor The operation of a reactor using 2.25Cr-1Mo-V steel, which is one of the Cr-Mo-V steels, is classified into start-up, steady-state operation, and shutdown, and the temperature, stress, and diffusible hydrogen distribution occurring during the operation of the reactor are simulated based on the relationship between temperature and pressure over time during these operations. The operating conditions of the 2.25Cr-1Mo-V steel reactor targeted in this embodiment, and the relationship between temperature and pressure over start-up and shutdown times are shown in Table 1 and Figure 1, and the basic design structure is shown in Figure 2. In this embodiment, ASME (SA-336 Gr. F22V) is used as the 2.25Cr-1Mo-V steel for the reactor material.

[0011] The basic structure of the reactor 1 will be described with reference to FIG. The reactor 1 has a shell 2 with a thickness of 228 mm excluding the WOL, a top head 2A above, a bottom head 2B below, and a skirt 2C. The shell and head are overlaid with 3.0 mm thick W / TP309L and 3.5 mm thick TP.347. The distance H between the top tangent line and the bottom tangent level is 34000 mm. The sidewall of the shell 2 is provided with a plurality of 6″ ID quench nozzles 3 and a plurality of 6″ ID catalyst drain nozzles 4 .

[0012] [Table 1]

[0013] The start-up period was 100 hours, the shutdown period was 57 hours, and the steady-state operation period was 17,363 hours, which was calculated by subtracting these periods from the two-year operation period per cycle.The reactor is composed of a SA-336M-F22V steel forged shell ring with an inner diameter of 4,740 mm and a wall thickness of 228 mm, which is a pressure-resistant member, and its inner surface is coated with a two-layer austenitic stainless steel overlay (hereinafter referred to as WOL) of 3 mm TP.309L and 3.5 mm TP.347 as a corrosion-resistant member.

[0014] In Figure 1, the temperature data versus time is taken as the temperature of the process fluid in the reactor, and a non-steady heat conduction analysis is performed by applying boundary conditions to the inner surface of the reactor to obtain the temperature distribution in the reactor wall. A structural analysis is performed using this temperature distribution as a thermal load to obtain the thermal stress distribution. Furthermore, a structural analysis is performed using the pressure data versus time in Figure 1 as the internal pressure and applying load conditions to the inner surface of the reactor to obtain the stress distribution throughout the wall thickness of the reactor while it is in service (the above analysis is performed, for example, using the methods of ASME Sec.VIII, Div.2 or Div.3).

[0015] The diffusible hydrogen distribution that occurs in the wall while the reactor is in service can be simulated by calculating the unsteady heat conduction equation, in which the hydrogen concentration C1 on the reactor inner surface given by the following equation (1) according to Sievert's law, the hydrogen concentration C3 on the 2.25Cr-1Mo-V steel side at the boundary between the pressure-resistant and corrosion-resistant components given by equation (2), and the hydrogen concentration C2 on the stainless steel WOL side given by equation (3) are each regarded as temperatures.

[0016] Based on the simulation results, the time history response of the hydrogen concentration on the inner surface of the reactor, the boundary WOL side, and the 2.25Cr-1Mo-V side according to the relationship between temperature and pressure versus time in Fig. 1 is shown in Fig. 3. Note that the hydrogen concentration C4 on the outer surface of the reactor is 0 ppm.

[0017]

number

[0018] Here, from Ahrrenius’ theorem, the hydrogen solubility K in 2.25Cr-1Mo-V steel and stainless steel WOL is b (T) and K C (T) and hydrogen diffusion coefficient D b (T) and D c (T) is given by equations (4) and (5), respectively. T and K are temperature and P and Pa are internal pressure. The temperature and pressure data against time in Figure 1 are given. L b and L c is the wall thickness, m, of 2.25Cr-1Mo-V steel and stainless steel WOL, L b =228e-3m, L c = 7.5e-3m, and 1 molH / m 3 =0.128wppmH.

[0019]

number

[0020] where the parameters K0, ΔH, D0 and E a The ideal gas constant R is shown in Table 2, and R = 8.314 J / (molK). In the hydrogen diffusion simulation, the specific heat C p Here, λ is the thermal conductivity, W / m o C), ρ is density, kg / m 3 It is.

[0021]

number

[0022] [Table 2]

[0023] The basic dimensions, components and materials of the 6”ID quench nozzle 4 that was the subject of the unsteady heat conduction analysis, stress analysis and hydrogen diffusion simulation are shown in Figure 4. The design structure satisfies the requirements of Part 4, Design by Rule, applying the allowable stress value of 199.8MPa at the design temperature of 454℃ for SA-336M-F22V steel in ASME Sec.VIII, Div.2.

[0024] FIG. 4 is a cross-sectional view showing the 6" ID quench nozzle 4. The 6" ID quench nozzle 4 has an outer diameter of 235 mm and protrudes from the shell 2. It has a quench pipe 40 that communicates with the inside of the shell in its central axial direction and has a flange with a width of 118 mm at its tip. The outer layer of the quench pipe 40 has a multi-layer shield 41. An insulation 43 is provided on the outer periphery of the flange via a gasket 42, and a reducing flange 44 is provided at the tip of the insulation 43.

[0025] The specifications and boundary conditions of the analysis model (unsteady heat conduction analysis / hydrogen diffusion simulation) are shown in Fig. 5, and the material properties applied to them are shown in Table 3.

[0026] [Table 3]

[0027] The data of the temperature of the process fluid in the unsteady heat conduction analysis, the hydrogen concentration in the hydrogen diffusion simulation, and the internal pressure in the structural analysis are given to Boundary1 in Fig. 5, which shows the inner surface of the reactor. The 2.25Cr-1Mo-V steel directly below the stainless steel WOL on the inner surface of the nozzle welded joint shown in Fig. 5 was selected as the analysis point. The temperature and diffusible hydrogen histories are shown in Fig. 6 as the unsteady heat conduction analysis results and the hydrogen diffusion simulation results, respectively, with the solid and dashed lines, and the histories of the circumferential and axial stress components, which are the stresses in the direction of crack opening, are shown in Fig. 7 as the structural analysis results, respectively. The maximum values ​​of the temperature and diffusible hydrogen content were 425°C and 5.20 ppm, respectively, during the steady operation period. Both the circumferential and axial stress components showed a stress increase 11 hours after the start of shutdown in the high temperature range of 400°C, where hydrogen solubility is high, and reached a maximum value of 194 MPa and 116 MPa, respectively, and then decreased for 2 hours, but the stress increase continued for 10 hours again. Therefore, it was confirmed that the rising load test was appropriate for reproducing the shutdown in the hydrogen embrittlement test.

[0028] 2. Reflection of reactor material properties in hydrogen embrittlement tests The material properties of the 2.25Cr-1Mo-V steel reactor are reflected in the test materials for the hydrogen embrittlement test. As shown in Figure 8, the decrease in tensile strength at the design temperature of 454°C due to the holding time at the PWHT temperature of 705°C for 2.25Cr-1Mo-V steel is shown in Figure 8. When manufacturing the reactor, the Min. PWHT is 8 hours at 705°C, and the Max. PWHT is 32 hours, with four PWHTs guaranteed, including repairs during service.

[0029] On the other hand, the transition temperature, which indicates the absorbed energy of 55 J in the Charpy impact test, shifts to the higher temperature side in the case of materials that are highly susceptible to temper embrittlement after step cooling, which is an accelerated embrittlement heat treatment. Figure 9 shows the results of a comparison between 2.25Cr-1Mo steel and 2.25Cr-1Mo-V steel. For 2.25Cr-1Mo-V steel, the shift in transition temperature due to step cooling is less than 10°C for the base metal, weld metal, and HAZ, confirming that the degree of temper embrittlement during service is low. Therefore, the level of toughness that may occur during reactor production is reflected in the test material.

[0030] The base material of the welded test plate was given the same tempering conditions as the shell of the actual reactor, and the thickness and groove shape of the shell ring circumferential weld joint of the actual reactor were reproduced on the welded test plate, and the test material was produced under the same welding conditions as the actual reactor. The groove shape and welding conditions of the welded test plate, the chemical composition of the weld metal at half its thickness, and the tensile test results and Charpy impact properties of the weld metal with PWHT temperatures of 705°C and 680°C are shown in Figure 10, Tables 4, 5, 6, and Figure 11, respectively. Table 5 shows the chemical composition of the welded part of the test material for the reactor material ASME material (SA-336 Gr.F22V).

[0031] [Table 4]

[0032] [Table 5]

[0033] [Table 6]

[0034] Whereas the weld metal with a PWHT temperature of 705°C for 8 hours met the SA-336M F22V specifications for the base metal, at 680°C the strength rose above the specifications on the high strength side, resulting in a decrease in toughness. The fracture appearance transition temperature (FATT) was -42°C for the weld metal with a PWHT temperature of 705°C and -16°C for the weld metal with a PWHT temperature of 680°C. Furthermore, the position of the notch relative to the toughness at 20°C did not affect the weld metal with a PWHT temperature of 705°C, but did affect the weld metal with a PWHT temperature of 680°C, resulting in a decrease in toughness at 1 / 4 width of the weld metal compared to the center. Based on these results, a high-toughness weld that was subjected to PWHT conditions of 705°C for 8 hours and with a notch position at the center of the weld metal, and a low-toughness weld that was subjected to PWHT conditions of 680°C for 8 hours and with a notch position at 1 / 4 the width of the weld metal were reproduced in the 1T-C(T) test pieces used in the rising load test.

[0035] 4. Reflection of phenomena occurring during reactor operation in hydrogen embrittlement tests The relationship between the phenomena occurring during reactor operation and the hydrogen embrittlement test conditions is shown in Table 7. The phenomena occurring during reactor operation are classified into temperatures, stresses, and diffusible hydrogen contents that occur during start-up, steady-state operation, and shutdown, and these conditions are reflected in the hydrogen embrittlement test to evaluate the hydrogen-assisted crack propagation initiation load and its propagation behavior.

[0036] [Table 7]

[0037] Since the stress generated in the reactor increases during both the start-up and shutdown periods, the rising load test is used to reproduce the hydrogen-assisted crack propagation behavior when the load increases. In addition, during steady-state operation, the holding load test is used to reproduce the crack propagation behavior under a constant load, in which the hydrogen-assisted crack propagation initiation load obtained from the rising load test is maintained.

[0038] 4.1 Reproduction of diffusible hydrogen in steel Figure 12 shows a schematic diagram of the IHAC progression in a reactor. During steady-state operation, the reactor reaches a high temperature of 400°C, and the solubility of atomic hydrogen in both austenitic stainless steel and Cr-Mo, which are the materials used, increases, so hydrogen is charged into the steel. When the steel is cooled to room temperature during shutdown, the atomic hydrogen in the steel becomes supersaturated, but degassing takes time, and diffusible hydrogen remains in the steel. As a result of the concentration of this diffusible hydrogen in the plastic region near the tip of the existing defect, this local high hydrogen concentration causes embrittlement, and in combination with stress concentration, microcracks are generated, which are thought to link together and appear as crack propagation. It is also thought that the supersaturated diffusible hydrogen accumulates in the intrinsic defect and becomes hydrogen gas, which acts as internal pressure to promote crack propagation. Such diffusible hydrogen in steel during reactor operation can be reproduced by exposing a test specimen to high-temperature, high-pressure hydrogen equivalent to that in the steady-state operation of the reactor in an autoclave and immediately cooling it with water.

[0039] In order to reproduce the diffusible hydrogen in the steel during use in the reactor, the hydrogen charging conditions for the test specimens and the hydrogen dissipation from the test specimens were investigated. Figure 13 shows the relationship between the time left in air at room temperature after hydrogen charging and the hydrogen content for hydrogen charging pressure and time of 25 MPa and 48 hours at temperatures of 375°C, 450°C, and 500°C, respectively. Since hydrogen dissipation was lowest at a hydrogen charging temperature of 450°C, and because the maximum amount of diffusible hydrogen generated in the pressure-resistant part during use in the reactor (5.20 ppm) was sufficiently exceeded even after leaving the specimen in air at room temperature for 9 hours, this condition was selected as the hydrogen charging condition.

[0040] Furthermore, after hydrogen charging, the test coupon was left in the atmosphere at 150°C, and the results of the investigation of the residual hydrogen content with analytical samples taken from the surface and center are shown in Figure 14. Even if hydrogen dissipates from the test coupon during the test at 150°C, it is expected that hydrogen exceeding the maximum diffusible hydrogen of 5.20 ppm that occurs in the pressure-resistant part during reactor operation will remain in the test coupon if the test is left for about 5 hours. In fact, after hydrogen charging the 1T-C(T) test coupon and conducting a rising load test in the atmosphere at 150°C, the hydrogen analysis results of samples taken at both ends and the center of the crack tip, and at a position half the distance from the crack tip to the bottom of the test coupon, and the thermal history of the analyzed samples are shown in Figure 15. It was confirmed that if the rising load test at 150°C is conducted within 4 hours, the hydrogen content in the crack propagation path of the hydrogen-charged 1T-C(T) test coupon can be reproduced to be equal to or greater than that in the pressure-resistant part of a reactor in service.

[0041] 4.2.Reproduction of dynamic load on a surface crack-like defect in high-pressure hydrogen gas Figure 16 shows a schematic diagram of HEAC propagation in the reactor. When the new surface at the tip of the surface crack is exposed to H2 gas with increasing load, H2 molecules decompose and atomic hydrogen is generated at the crack tip. This increasing stress field at the tip of the surface crack affects the balance between H2 molecules and atomic hydrogen, resulting in extremely high hydrogen concentration near the tip of the surface crack due to hydrogen trapping, which is thought to cause localized hydrogen embrittlement and manifest as HEAC propagation. Therefore, we applied an increasing load to a 1T-C(T) specimen in an autoclave in 20 MPa H2 gas, which is equivalent to the operating pressure of the reactor, and maintained the hydrogen-assisted crack propagation initiation load obtained from the load increase to reproduce the generation of a new surface at the tip of the surface crack and its exposure to H2 gas.

[0042] 4.3. Definition of hydrogen-assisted crack propagation initiation load As shown in Figures 12 and 16, respectively, both the IHAC and HEAC in the reactor are promoted by the plastic zone near the crack tip, so the hydrogen-assisted crack propagation initiation load is more sensitive to load increase than in the static state. In the rising load test used to define the hydrogen-assisted crack propagation initiation load in this study, the load was increased at a positive loading rate dK / dt, and the deviation was determined as the point at which the load-displacement curve obtained for the hydrogen-charged specimen, the test in 20 MPa H2 gas, or both, began to deviate from the curve obtained for the standard test in air using a specimen not charged with hydrogen. The value of K in equation (7) from ASTM E 1820 calculated using the loads P and N at this deviation point was taken as K. IH As the K for a crack-like defect occurring in the reactor, defined by equation (8), I By comparing this value with the actual value, the criterion for the onset of hydrogen-assisted cracking propagation during reactor operation is defined.

[0043]

number

[0044] where a is the crack length of the specimen, m, W is the specimen width, m, B, B N is the specimen thickness, m. Similarly, hydrogen-assisted crack propagation is governed by the concentration of atomic hydrogen in the plastic zone near the crack tip. IH The value of increases monotonically with the loading rate, and this phenomenon is also seen on the fracture surface. As the loading rate increases, the effect of hydrogen embrittlement on the fracture decreases. Therefore, in order to measure the load at which hydrogen-assisted crack propagation begins more conservatively, a slower loading rate is required. Therefore, in the rising load test in this study, the crosshead speed was set to 0.01 mm / min, which is the lower limit of the device, and the loading rate was dK / dt = 1.3 × 10 -2 The test was carried out at MPa√m / s√m / s.

[0045] 5. Identification of hydrogen-assisted crack propagation behavior According to the test conditions in Table 7, hydrogen embrittlement tests were carried out by reproducing the phenomena occurring during reactor start-up, steady-state operation and shutdown.

[0046] 5.1. Hydrogen-assisted crack propagation behavior during start-up Whether IHAC overlaps with HEAC was evaluated by comparing the load-displacement curves of specimens used in rising load tests in 20 MPa H2 gas when hydrogen was charged and when not. The load-crosshead displacement curves of both tests were plotted with the difference in load line displacement from the standard test added to determine the hydrogen-assisted crack propagation initiation load, and the corresponding fracture surface was evaluated.

[0047] The graphs of evaluation at room temperature for high and low toughness welds and the results of SEM fracture surface observations are shown in Figures 17 and 18, and Figures 19 and 20, respectively. Both high and low toughness welds showed pseudo-cleavage crack growth regardless of whether hydrogen was charged or not. In the case of high toughness welds, IH Since there was almost no difference in the values ​​of IHAC due to diffusible hydrogen in the steel, it is judged that the superposition effect of IHAC on the HEAC propagation behavior with increasing load does not occur. In the case of low toughness welds, regardless of whether hydrogen charging was performed or not, the pseudo-cleavage crack propagation led to fracture when the load increased.

[0048] Since the fracture initiation point was reached with almost no increase in the crosshead displacement during these load increases, it was assumed that the pseudo-cleavage crack had progressed to cleavage fracture, and the K value that gives the limit was set as K IC-H These K IC-H The value of is taken as the fracture toughness value affected by hydrogen embrittlement, and the fracture toughness value K obtained from the standard test is IC It is considered that the value is lower than the value obtained without hydrogen charging. IH and K. IC-HThe value of is lower than that obtained when hydrogen was charged, suggesting that not only is there no superposition of IHAC, but also that atomic hydrogen that penetrates from the newly formed surface as the load increases brings about a very high hydrogen concentration near the crack tip before the diffusible hydrogen in the steel does, causing localized hydrogen embrittlement. A schematic diagram of these HEAC propagation behaviors is shown in Figure 21.

[0049] The graphs and SEM fracture surface observation results for high and low toughness welds evaluated at 150°C are shown in Figs. 22 and 23, and Figs. 24 and 25, respectively. Both high and low toughness welds exhibited ductile crack growth regardless of whether hydrogen was charged or not. IH Since there is almost no difference in the values ​​of IHAC and IHAC due to diffusible hydrogen in steel, it is judged that the superposition effect of IHAC on the HEAC growth behavior with increasing load does not occur. IH Since there was no difference in the values ​​of , it is concluded that the effect of high or low toughness on the HEAC propagation behavior has disappeared. A schematic diagram of these HEAC propagation behaviors is shown in Figure 26. Since IHAC does not overlap at either room temperature or 150°C, it is believed that atomic hydrogen penetrates from the newly formed surface as the load increases, causing microscopic cracks to occur closer to the crack tip in the HEAC than in the IHAC.

[0050] 5.2. Hydrogen-assisted crack propagation behavior during steady-state operation The propagation of HEAC under constant load conditions was evaluated by a combination of rising and holding load tests in 20 MPa H2 gas at room temperature using hydrogen-charged specimens. For high and low toughness welds, the load was increased to the pseudo-cleavage crack propagation initiation load in the rising load test, and the load was held for up to 10 days in the holding load test. In both cases, crack propagation continued after the pseudo-cleavage crack propagation initiation load was held, and the load-load line displacement curves and the load line displacement history for the pseudo-cleavage crack propagation initiation load for 10 days are shown in Figures 27 and 28 for the high toughness welds and Figures 29 and 30 for the low toughness welds.

[0051] In addition, as shown in Fig. 31, the crack length that is thought to have grown during the load holding period was measured based on the fracture surface SEM images obtained in this test, and the correlation between the load line displacement increment and the crack length increment was calculated. The results are shown in Fig. 32. Focusing on the fact that the change in K value during quasi-cleavage crack growth depends on the crack length increment, a crack growth curve and the associated K value history were obtained from the correlation between the load line displacement increment and crack length increment in Figure 32, and the crack growth rate da / dt and load rate dK / dt were numerically introduced by time differentiation. The Paris law was applied to the relationship between da / dt and K, and the effect of dK / dt on this relationship was used to numerically model the growth of a quasi-cleavage crack during load holding as the crack growth rate, as shown in equation (9) for high toughness welds and equation (10) for low toughness welds.

[0052]

number

[0053] In addition, since the crack growth rate at room temperature experimentally shows a linear temperature dependence with respect to the inverse of the temperature, 1 / T, the temperature dependence of the crack growth rate can be modeled using equation (11).

[0054]

number

[0055] A comparison of the crack growth rate models for high and low toughness welds is shown in Figure 33, and a comparison of the crack size increments calculated from these numerical models, assuming a steady-state operation period of two years per reactor cycle, is shown in Table 8.

[0056] [Table 8]

[0057] If the loading rate dK / dt is the same, that is, the contact rate of H2 gas to the newly formed surface is the same, the crack growth rate is significantly faster in low toughness welds than in high toughness welds. In addition, there are dK / dt conditions where the crack growth rate approaches a plateau in both high and low toughness welds, and compared to high toughness welds, low toughness welds initiate at a higher dK / dt and their crack growth rate is also significantly faster. In an actual reactor, the cracking of the inner stainless steel WOL that has progressed to the pressure-resistant components and the cracking from the peeling part of the inner stainless steel WOL to the pressure-resistant components correspond to the HEAC obtained in this study. The crack size increments from the start of the pseudo-cleavage crack growth calculated for the steady-state operation period of the reactor for two years are all at a level that can be detected by nondestructive testing during shutdown maintenance. Therefore, the validity of the growth of crack-like defects predicted by these crack growth rate models can be compared and verified by follow-up surveys at each shutdown maintenance.

[0058] 5.3. Hydrogen-assisted crack propagation behavior during shutdown The presence or absence of IHAC progression at 150°C and room temperature was evaluated using the load-displacement curves of a rising load test in air using hydrogen-charged specimens. However, based on the results of the study in section 4.1, it is predicted that the crack progression path of the hydrogen-charged 1T-C(T) specimen in the rising load test at 150°C contains hydrogen equal to or greater than the diffusible hydrogen at 425°C during reactor shutdown. The load at which hydrogen-assisted cracking begins to propagate was defined using a graph that added the difference in load line displacement from the standard test for each temperature to the load-crosshead displacement curves obtained from the tests at 150°C and room temperature, and the corresponding fracture surface was observed.

[0059] Graphs of evaluations of high-toughness welds at 150°C and room temperature and the results of SEM fracture surface observations are shown in Figures 34 and 35, respectively. IHAC did not occur at room temperature, only N2 occurred at 150°C. The minimum value of residual hydrogen measured after the test was 6.66 ppm, which exceeds the simulation result of 5.20 ppm, so it is determined that the test itself is valid and that IHAC appears at 150°C, where the solubility of atomic hydrogen in steel is high.

[0060] The graphs and SEM fracture surface observation results for low-toughness welds at 150°C and room temperature are shown in Figures 36 and 37, respectively. IHAC did not occur at room temperature, but did occur at 150°C. As with high-toughness welds, it is believed that IHAC appears at 150°C, where the solubility of atomic hydrogen in steel is high.

[0061] High and low toughness welds IH Since there was no difference in the values ​​of , it can be determined that the effect of high and low toughness on the IHAC progression behavior has disappeared. A schematic diagram of these IHAC progression behaviors is shown in Figure 38. The hydrogen embrittlement resistance characteristics at room temperature are thought to be due to hydrogen being trapped at the interface between the vanadium carbide and the matrix, and are manifested as the characteristic that IHAC progression does not occur. The reason that IHAC appears at 150°C, where the solubility of atomic hydrogen in steel is high, is thought to be because the hydrogen trapped by the vanadium carbide diffuses again.

[0062] In order to set the minimum pressurization temperature and evaluate the life of 2.25Cr-1Mo-V steel reactors for oil refineries, fracture mechanics properties for HEAC obtained from rising load tests in high-pressure H2 gas equivalent to the operating pressure of the reactor are required.

[0063] K obtained in a rising load test in 1.20MPa H2 gas IH The value of HEAC is dominant and IHAC is hardly overlapped. The behavior shows pseudo-cleavage crack growth at room temperature and ductile crack growth at 150℃. For low toughness welds at room temperature, the fracture toughness value K due to hydrogen embrittlement is IC-H This suggests a transition to cleavage fracture.

[0064] In the holding load test under the pseudo-cleavage crack propagation initiation load in 2.20 MPa H2 gas, the crack propagation continued, and the crack propagation rate was suggested to have a plateau characteristic in the relationship between da / dt and K and a dependence on dK / dt.

[0065] 3. IHAC appeared in a rising load test at 150℃ in air using hydrogen-charged material, but did not appear at room temperature. This is thought to be because the hydrogen trapped by VC carbide diffuses again at high temperatures where the solubility of atomic hydrogen in steel increases.

[0066] In the above embodiment, the hydrogen embrittlement was evaluated for the reactor using 2.25Cr-1Mo-V steel, but the same evaluation can be performed for Cr-Mo-V steels having compositions other than the above. Furthermore, the applications are not limited to reactors, but can also be used for reactor peripheral equipment and piping, etc., where startup operation, steady state operation, and shutdown operation are performed as one cycle in a hydrogen environment. EXAMPLES

[0067] (Step 1) For the test materials shown in the embodiment, the basic design structure of the actual reactor is the same as that shown in FIG. 2 of the embodiment, and the operating conditions are shown in Table 1. The reactor material is ASME SA-336 Grade F22V (forged material) as shown in the embodiment. The specific composition (mass %; balance Fe and other unavoidable impurities) is shown below. C: 0.11~0.15 Mn: 0.30~0.60 P:≦0.015 S:≦0.010 Si:≦0.10 Ni:≦0.25 Cr: 2.00~2.50 Mo: 0.90~1.10 V: 0.25~0.35 Nb:≦0.07 B:≦0.0020 Ti:≦0.030 Cu:≦0.20 Ca:≦0.015

[0068] As in the above embodiment, the relationship between the internal fluid temperature and the internal pressure with respect to the start-up and shut-down times of the 2.25Cr-1Mo-V steel reactor shown in FIG. 12 was used as the reactor operating conditions for the simulation.

[0069] (Step 2) Furthermore, the time history responses of the hydrogen concentration on the inner surface of the reactor, the boundary WOL side, and the Cr-Mo side in response to the start-up and shutdown of the reactor, shown in Fig. 13, were set in the hydrogen diffusion simulation.

[0070] In order to model the crack-like defect to be evaluated, the axial semi-elliptical crack model on the inner side of the 6” ID Quench Nozzle welded joint shown in Fig. 5 was set as the analysis point, and the crack depth a was set to 18 mm and the crack depth l was set to 2c = 108 mm.

[0071] The calculated stress state and diffused hydrogen content of the crack-like defect model for the reactor operating conditions are shown in Figure 5.

[0072] (Step 3) The simulated time history responses of temperature and hydrogen accumulation (thermal conduction analysis results and hydrogen diffusion simulation results) for reactor startup and shutdown are shown in Figure 6. In the calculation results, the upper limit of hydrogen concentration reaches 5.20 ppm, and the steady-state operating temperature is set to 425°C. Then, the time history response of the crack opening stress components for the start-up and shutdown of the reactor is obtained by structural analysis, as shown in Fig. 7. The circumferential stress is 187 MPa during steady operation and increases to 194 MPa during shutdown. The axial stress is 110 MPa during steady state operation and rises to 116 MPa during shutdown.

[0073] (Step 4) The hydrogen embrittlement material properties are predicted from the Charpy impact properties at the time of reactor fabrication and the hydrogen content during service. The relationship between load and crosshead displacement is shown in Figure 39. The criteria for evaluating the propagation of hydrogen-assisted cracks are shown below. The rapid fracture criterion for hydrogen-assisted crack propagation is given below.

[0074]

number

[0075]

number

[0076] (Step 5) Next, the hydrogen-assisted cracking propagation tendency is evaluated according to the reactor operating conditions. In the hydrogen-assisted crack propagation evaluation criterion, hydrogen-assisted cracking propagates in the crack depth direction just before the start of steady-state operation (load increase and hold). The hydrogen-assisted crack propagation evaluation results (time history response of stress intensity factors for reactor startup and shutdown) are shown in Figure 40.

[0077] (Step 6) A determination is made as to whether the reactor can be continued to be operated in the event of hydrogen embrittlement accompanied by the progression of hydrogen-assisted cracking. The criterion for determining whether or not the reactor can be continued to operate due to hydrogen embrittlement is K. mat =K IC-H Since continued operation is deemed acceptable, a remaining life assessment will be carried out using crack propagation analysis. mat is the lowest limit of the fracture toughness value when variation is taken into consideration. The results of the determination of whether or not the reactor can be continued to be operated in the case of hydrogen embrittlement accompanied by the progression of hydrogen-assisted cracking are shown in Figure 41. If it is in the acceptable region of the fracture assessment diagram (FAD), it is possible to continue operation, but if it is in the unacceptable region, it is not possible to continue operation.

[0078] (Step 7) The remaining life of the reactor will be evaluated by hydrogen-assisted crack propagation analysis. The depth of the cracks that occurred in the test material was measured using SEM images (Figure 42; fracture surface observation results at 200x using FE-SEM), and the correlation between the crack depth increment and the load line displacement increment was determined. The correlation was performed using the following formula: The obtained correlation diagram is shown in FIG. From Fig. 43, the pseudo-cleavage crack propagation start (K IH value) to cleavage fracture (K IC-H ) the subcritical crack growth rate can be predicted.

[0079]

number

[0080] Figure 44 shows the results of hydrogen-assisted crack growth analysis (crack growth curve and stress intensity factor distribution) for repeated reactor operation.

[0081] Next, the crack growth evaluation points according to da / dt are plotted on the failure assessment diagram (FAD) (Figure 45). As is clear from this figure, the seventh steady-state operation is within the acceptable region, but the seventh steady-state operation reaches the unacceptable region and is judged to be unacceptable (rapid destruction). In other words, the remaining life is allowed for six operations (12 years = 2 years / operation x 6 service periods).

[0082] For the axial semi-elliptical crack on the inner side of the 6" IDQuench Nozzle welded joint, Initial crack depth: α0 = 18 mm Initial crack length: l0 = 2c0 = 108 mm Critical crack depth: α c =56.40mm Critical crack length: l c =2c c =108mm. Since it is predicted that no crack surface growth will occur during the allowable period of continued operation, it is necessary to continuously monitor the growth of the crack in the depth direction by ultrasonic testing from the outer surface of the reactor. [Explanation of symbols]

[0083] 1. Reactor 2A Top Head 2B Bottom Head 2C Shell Bottom Head 3 6” ID Quench Nozzle 4 6” ID Catalyst Drain Nozzle

Claims

1. A method for evaluating hydrogen embrittlement of a component having Cr-Mo-V steel, which is operated in a cycle of startup, normal operation and shutdown and is used in a hydrogen environment, comprising: obtaining information on the structure and physical properties of the material constituting the member, and information on the process temperature, process pressure, and hydrogen concentration under each operating environment according to the operation time at the start-up, the steady operation, and the shutdown, and the elapsed operation time; A step of simulating the phenomena of temperature, stress distribution through the thickness, and diffusible hydrogen distribution occurring in the component during operation based on each of the information; a real-machine characteristic imparting step of subjecting a specimen reflecting the component to manufacturing conditions executed in manufacturing the component to impart material characteristics of a real-machine material; An evaluation step of classifying phenomena occurring during service into temperatures, stresses, and diffusible hydrogen contents occurring during start-up, steady-state operation, and shutdown, and evaluating the progress of hydrogen-assisted cracking by reflecting these conditions in the test material; The hydrogen embrittlement evaluation method for Cr-Mo-V steel has the following features.

2. 2. The method for evaluating hydrogen embrittlement of Cr-Mo-V steel according to claim 1, wherein, in the actual equipment characteristic imparting step, the test material is given the tempering conditions executed for the actual Cr-Mo-V steel, the shape of a welded joint is reproduced in the test material, and welding conditions including post-weld heat treatment are given.

3. 3. The method for evaluating hydrogen embrittlement of Cr-Mo-V steel according to claim 2, wherein the test materials are a high-toughness welded portion having a notch position at the center of the weld metal, assuming that the post-weld heat treatment temperature in an actual machine is on the high temperature side, and a low-toughness welded portion having a notch position that is 1 / 4 the width of the weld metal, assuming that the post-weld heat treatment temperature in an actual machine is on the low temperature side.

4. 4. The hydrogen embrittlement evaluation method for Cr-Mo-V steel according to claim 1, wherein in the evaluation step, the test material is exposed to a hydrogen environment simulating the temperature and hydrogen pressure during operation of an actual machine in service, and diffusible hydrogen is reproduced by hydrogen charging, and for startup and shutdown, a rising load test is used to reproduce the hydrogen-assisted crack propagation behavior when the load increases, and for steady operation, a holding load test is used to reproduce the crack propagation behavior under a constant load, in which the hydrogen-assisted crack propagation initiation load obtained from the rising load test is held.

5. In the rising load test, among the test specimens to which the characteristics of the actual machine were imparted, the load-displacement change was measured for test specimens that were hydrogen-charged or placed in a hydrogen environment without being hydrogen-charged, or test specimens that were hydrogen-charged and placed in air, and test specimens that were not hydrogen-charged and placed in air. The K value at which the measurement results of the test specimens that were hydrogen-charged or placed in a hydrogen environment without being hydrogen-charged, or test specimens that were hydrogen-charged and placed in air start to deviate from the measurement results of the test specimens that were not hydrogen-charged and placed in air was calculated as K. ＩＨ The K value, which indicates the point at which fracture occurs without an increase in displacement when the load is increased, is defined as K. ＩC-Ｈ and K is the K value calculated for a detected crack-like defect or a hypothetical crack-like defect under actual load conditions. Ｉ Compared with the value, K Ｉ The value is K ＩＨ When the value is equal to or greater than this, it is regarded as the criterion for the start of hydrogen-assisted crack propagation. Ｉ The value is K ＩC-Ｈ The method for evaluating hydrogen embrittlement of a Cr-Mo-V steel according to any one of claims 1 to 3, wherein a value equal to or greater than this value is used as a criterion for determining whether hydrogen-assisted cracking progresses to rapid fracture.

6. 6. The method for evaluating hydrogen embrittlement of a Cr-Mo-V steel according to claim 5, wherein in the steady-state operation test, the load is increased to a pseudo-cleavage crack propagation start load, which is a hydrogen-assisted crack propagation start load, by a rising load test, and then a holding test is performed.

7. 4. The method for evaluating hydrogen embrittlement of a Cr-Mo-V steel according to claim 1, further comprising the steps of: measuring a crack length of a test material obtained by testing the test material during a pseudo-cleavage crack propagation period after hydrogen-assisted crack propagation starts; determining a correlation between a load line displacement increment and a crack depth increment; predicting a crack propagation rate from the correlation; and estimating a time when a K value giving the crack propagation rate reaches a hydrogen embrittlement fracture toughness value.

8. The method for evaluating hydrogen embrittlement of a Cr-Mo-V steel according to any one of claims 1 to 3, wherein the Cr-Mo-V steel is a 2.25Cr-1Mo-V steel and has a composition specified in ASME SA-336 Grade F22V or ASME SA-542 Type D, Class 4a.

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