Effects of normalizing temperature on microstructure and high-temperature properties of 2.25Cr-1Mo-0.25V bainitic steel

In 2.2Cr-1Mo-0.25V bainitic steels, the effect of the initial microstructure on carbide precipitation has not yet been elucidated. Therefore, in this work, the effect of normalizing temperature on the microstructure evolution and corresponding high-temperature strength change of bainitic steels was investigated by transmission electron microscopy (TEM) and scanning electron microscopy (SEM) using extra-thick hydrogenation reactor steel as the assessed material, and the precipitation strengthening mechanism was discussed. When the normalizing temperature increased from 900 °C to 990 °C, the high-temperature yield strength increased from 451 to 475 MPa and the average absorbed Charpy impact energy at −30 °C decreased from 93 to 41 J. Therefore, the optimal normalizing temperature was 960 °C. The volume fraction of MC carbides with almost no change in average diameter increased from 0.21% to 0.44%, and the dispersed strengthening of carbide was significantly enhanced. The increase in normalizing temperature increased the solid solution of alloying elements within the bainite matrix, which increased the precipitation of MC carbides during the tempering process. In addition, the lamellar M-A constituents in the initial microstructure decomposed to form MC carbides.


Introduction
In the oil-refining sector, hydrogenation reactors must run for extended periods of time under hightemperature, high-pressure, and hydrogen conditions [1]. Low alloy 2.25Cr-1Mo-0.25V steel, the material of hydrogenation reactors, requires outstanding raised temperature performance and resistance to oxidation for safe operation [2][3][4]. The size and cross-section of pressure vessels can also be expanded to achieve improved processing efficiency, due to an increase in energy consumption [5]. The traditional heat treatment procedure for creating 180 mm thick low alloy 2.25Cr-1Mo-0.25V steel involves normalizing and tempering [6]. In general, during normalization, distinct microstructures will form from the surface of the component to its center, due to the large cross-temperature of the section's gradient [7]. The middle portion of the component, which generally has a weak connection with the mechanical characteristics, typically contains granular bainite because of the slower cooling rate [8]. The initial phase in the heat treatment process for extra thick 2.25Cr-1Mo-0.25V steel plates consists of the normalizing process [9]. In general, determining the connection between microstructure and normalizing temperature is straightforward. However, understanding exactly how the normalized structure affects the mechanical material properties because of the complicated microstructure evolution during tempering remains a challenge. Therefore, to increase the dependability of hydrogen reactors, it is critical to evaluate the microstructure and mechanical properties of extra thick 2.25Cr-1Mo-0.25V steel plates.
Numerous factors in normalized microstructures, including grain size and bainite type, will be reflected as a result of the normalizing temperature [10]. The mechanical properties of 2.25Cr-1Mo-0.25V steel will exhibit considerable changes throughout the tempering process due to the decomposition of martensite-austenite (M-A) and the softening of bainite ferrite in the normalized microstructure, which will be impacted by the initial Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. normalized microstructure [11]. Jiang et al [12][13][14] discovered that the tempering temperature and time affected the precipitation behavior and coarsening of carbides in the tempering of 2.25Cr-1Mo-0.25V steels with a granular bainitic microstructure. Song et al [15][16][17] demonstrated that creep voids in 2.25Cr-1Mo-0.25V steel with granular bainite as the starting sprout accumulated at inclusions and coarsened carbides, inducing creep fracture. According to Li et al [18], the difference in strength between tempering the granular bainite microstructure at 650°C versus 700°C in 2.25Cr-1Mo-0.25V steel was mostly caused by the coarsening behavior of the carbides. Schönmaier et al [9,19,20] found that the type and quantity of carbides in 2.25Cr-1Mo-0.25V steel correlated with the time of the heat treatment process, with a substantial impact on the mechanical properties of the material. However, the aforementioned research was conducted for particular microstructures and did not account for the impact of normalizing microstructure variations on the mechanical properties.
Recent research has mainly focused on the correlation between the mechanical properties, temper brittleness, creep behavior, hydrogen embrittlement, and post-weld heat treatment and microstructural variations in 2.25Cr-1Mo-0.25V steel [3,[21][22][23][24][25]. Few studies have been conducted on this topic, despite the fact that the high-temperature tensile strength is a crucial mechanical indicator for the design and assessment of prohydrogen reactors. Song et al [15] examined the high-temperature tensile and creep behavior of 2.25Cr-1Mo-0.25V steel and weld metals. The weld metal outperformed the base metal in terms of performance due to fine grains, and both the base metal and weld metal experienced creep damage through second phase particle cavity sprouting and aggregation. Zhang et al [26] studied the microstructure and tensile properties of 9Cr-1Mo steel at high temperatures and discovered that tempering enhanced precipitation strengthening, which increased the high-temperature strength. Samant et al [27] studied the impact of thermomechanical treatment on the hightemperature tensile characteristics of 9Cr-1Mo heat-resistant steel, and showed that the high-temperature strength depended on the variability of the carbide diameter, area percentage, and number density. Although the behavior of carbide precipitation on Cr-Mo steels was clarified by these studies, it remains unknown how carbide precipitation affects the high-temperature mechanical properties of 2.25Cr-1Mo-0.25V steels with various normalized organizational structures after tempering.
In this study, we attempted to increase the high-temperature strength and fracture toughness by varying the normalizing temperature to obtain various normalizing microstructures. Tempering tests were conducted on specimens with various microstructures that were normalized. SEM and TEM analyses were used to analyze the microstructure evolution, including austenite grain size and M-A composition type, as well as the type, size, and distribution of precipitates. In addition, room temperature tensile, high temperature tensile, and −30°C Charpy impact tests were conducted to determine the mechanical properties. The link between the hightemperature mechanical properties and microstructure was discussed, along with the influence of normalizing temperature on the microstructure.

Experiments and methods
The chemical composition of the industrially manufactured steel plate was Fe-0.14C-2.3Cr-1Mo-0.25V-0.5Mn-Nb-Ti. The experimental steel ingots were forged into 180 mm-thick steel blocks. The steel blocks were heated to 1200°C at a heating rate of 350°C h −1 in a heat treatment furnace and maintained for 2 h for austenitizing and homogenizing. Two-stage hot rolling was carried out immediately by four-high hot mill (NEU, Liaoning, China) after heat preservation. In the first stage, the starting rolling temperature was 1050°C, and the final rolling temperature was 1000°C. After a certain time of holding the temperature at 880°C, the second stage of rolling was carried out. The final rolling temperature was 830°C, and then air cooling was carried out to room temperature. Finally, the billet was rolled into a 16 mm thick experimental steel plate. Then, a K-type thermocouple attached to a thermocouple thermometer was inserted and loaded into a heating furnace with the test material after drilling a hole in the middle of the thickness direction of the experimental steel, as shown in figure 1(a). The test material was then heated to 960°C using a heating furnace at a heating rate of 10°C min −1 , maintained for 40 min, and then cooled to room temperature by the air outside the furnace to acquire a cooling curve of the test material in air. The experimental steel, with a thickness of 16 mm, was cooled in air at a rate of roughly 0.5°C s −1 , according to the calculations ( figure 1(b)). According to the actual situation in the industrial field, the cooling speed of the 180 mm steel plate in the upper center of the roll quenching machine was about 0.3 ∼ 1°C s −1 , which was close to the cooling speed of small samples in the laboratory. Therefore, small samples could be used to simulate the core of industrial extra-thick plates. Normalizing and tempering were conducted on the experimental steel. Figure 1(c) depicts the procedure of heat treatment. After being normalized at various temperatures for 40 min, the experimental steel was cooled in air to room temperature. All of the normalized experimental steels were subsequently tempered for a total of 55 min at 730°C, and then air cooled to room temperature.
Standard cylindrical room-temperature tensile test samples with a gauge length of 25 mm and diameter of 5 mm, as well as high-temperature tensile test samples with gauge lengths of 25 mm and diameters of 6 mm, were processed in accordance with the GB/T 228.1-2010 and GB/T 228.2-2015 standards. The tests were carried out on an INSTRON-5582 tensile tester at constant crosshead speeds by an Instron Dynatup 9250 instrumented impact machine (Norwood, MA, USA), on which Charpy impact tests were conducted at −30°C. All of the samples were created as conventional Charpy V-notch specimens with dimensions of 10 × 10 × 55 mm.
The experimental steel was austenitized for 40 min at the corresponding normalizing temperature and then quenched in water to measure the austenite grain size. Subsequently, the prior austenite grains (PAG) of the normalized samples were viewed using optical microscopy and measured using the line intersection point method (ASTM E112) via Image-Pro Plus software. To study the microstructural characteristics, a fieldemission TEM and field-emission scanning electron microscopy (FE-SEM, ULTRA 55, Zeiss, Jena, Germany; FE-TEM, Tecnai G2 F20; FEI, Hillsboro, OR; USA) were utilized. The samples were prepared for SEM analysis by mechanically grinding, polishing, and chemically etching the samples with 4 vol.% natal alcohol solution. Then, the thin foil specimens were polished using a twin-jet technique at 20 V using a solution consisting of 10% perchloric acid and 90% ethanol. Energy-dispersive spectroscopy (EDS) was used to analyze the chemical makeup of the precipitates. The equilibrium phase diagram between the fraction and temperature of the phases was identified using thermo-calc software and the TCFE9 database in accordance with the chemical makeup of the 2.25Cr-1Mo-0.25V steels.  1020°C) experienced a steady increase from 451-505 MPa, and the third stage (1020°C-1050°C) underwent a minor decline in high-temperature strength. YS and UTS were at a maximum at room temperature when the normalizing temperature was 960°C, as shown in figure 2(b), and then they fell rapidly as the normalizing temperature rose to 990°C. Figure 2(c) shows that as the normalizing temperature increased from 870°C to 1050°C and the impact energy decreased from 111 J to 9.6 J. According to the GB/T 35012-2018 standard, the impact absorption energy at −30°C should be greater than 55 J, the high-temperature yield strength at 500°C should be greater than 325 MPa, and the room-temperature yield strength should be greater than 620 MPa. Our testing findings met the mechanical property standards, as the high experimental temperature (550°C) exceeded 500°C. The best normalizing temperature range, therefore, was found to be 900°C-960°C, and the 900°C-990°C temperature range was used in the studies discussed below.

Results
3.2. Normalized microstructure of 2.25Cr-1Mo-0.25V steel Figures 3(a)-(d) displays the shape and size of prior austenite grain (PAG) in the experimental steel, following normalization at 870, 930, 990, and 1050°C for 40 min. It was evident that the prior austenite grains expanded as the normalizing temperature increased. Figure 4 illustrates the average values of the PAG dimensions as a function of the normalizing temperature using statistics from image-pro plus software. The PAG growth stages could be categorized as slow and quick. In the slow growth stage (870°C to 930°C), the normalization temperature was lower, and the microstructure of the normalized sample was composed of coarse and fine mixed PAG. In the fast growth stage (930°C to 1050°C), the normalizing temperature at 930°C was the point at which the speed of PAG growth changed, and the homogeneous equiaxed PAG size increased significantly, from 9.68 to 43.9 μm. Figures 5 and 6 display the SEM and TEM images of the normalized microstructures, which were subjected to various normalizing temperatures. The entire prior austenite grain boundary was visible because bainite development adhered to the edge growth and displacive mechanisms. As the normalizing temperature increased, we found that the microstructure transformed from typical granular bainite to lath bainite. The morphology of M-A was one of the distinctions between lath bainite and granular bainite. Further TEM observations were made of the detailed M-A microscopic characteristics of the constituents. According to the high magnification image shown in figure 5(a), the martensite in the M-A constituents exhibited very fine twinning. Blocky and film M-A constituents with high density and entangled dislocations were also observed, as illustrated in figures 6(b), (c). This could be recognized as retained austenite based on the diffraction pattern at D2 of the membrane M-A island in figure 6(d). According to the crystallographic orientation relationship, bainitic ferrite was (111)/(011) and [011 - α, which was in agreement with the illustrated K-S orientation relationship. According to figure 5, the specimens normalized at 900°C had blocky M-A constituents that were randomly distributed inside the equiaxed bainite ferrite, while the specimens normalized at 930°C had film M-A islands sandwiched between the lath bainite ferrite, along with a few M-A constituents. After the normalizing temperature was raised, the volume percentage of the film M-A constituents increased.  identified as M 23 C 6 carbides. According to the EDS analysis results, the rod-like carbides shown in figure 8(b), ranging in size from 10 to 20 nm, could be recognized as V-rich MC carbides. Additionally, rod-like M 7 C 3 carbides larger than 100 nm were observed, which was in agreement with the findings of previous investigations [7]. Large spherical M 23 C 6 carbides were mostly localized at the grain boundaries, while elongated M 7 C 3  carbides and fine V-rich MC carbides were spread throughout the matrix, in accordance with the SEM and TEM images.
As the normalizing temperature was increased from 900°C to 990°C, the size and distribution of the precipitates changed. The average diameter, volume percentage, and number density of the carbides were calculated by image-pro plus software. At least 1600 M 23 C 6 and M 7 C 3 carbides were counted in the SEM images magnified by 8 x. The same method of counting the MC carbides was used for the TEM images, indicating that M 23 C 6 and M 7 C 3 had the lowest average diameters of around 130 nm and the highest number density of 2.82 μm −2 in the tempered specimens normalized at 900°C. When the normalizing temperature exceeded 930°C, the average diameter increased and the number density slightly decreased. The average diameter of the MC carbides nearly remained mostly unchanged when the normalizing temperature was increased from 900°C to 990°C, while the volume fraction significantly increased from 0.21% to 0.44% (table 1).

Effect of normalizing temperature on the microstructure evolution
Zeng et al [29] emphasized that NbC and VC carbides in high-strength low-alloy oil pipes and tube steels inhibited the growth of prior austenite grains during austenitization. According to Li et al [30], large-sized M 23 C 6 carbides were eliminated by austenitization at 1050°C in martensitic heat-resistant steels. In this study, undissolved carbides with diameters larger than 100 nm were observed in the normalized microstructure at 900°C , as shown in figure 9, in contrast to the aforementioned steels. Additionally, (Ti,V,Nb)C carbides were observed, as shown in figure 9(d). In the normalizing temperature from 900°C to 960°C, the amount of carbides inside the normalized microstructure decreased. According to the Zener pinning theory, the effect of austenite grain growth could be described by where P is the Zener pinning pressure, γ is the grain boundary energy, β is a dimensionless constant, and f and r denote the volume fraction and size of the carbides, respectively. The volume fraction of carbide decreased as a result of carbides dissolving into the austenite matrix as the normalizing temperature increased, and the pinning pressure for the growth of austenite grains also decreased. Consequently, the prior austenite grain sizes quickly increased in the 2.25Cr-1Mo-0.25V steel when the normalizing temperature exceeded 930°C. Figure 10(a) shows the thermodynamic calculation results obtained from preliminary analysis of the experimental steel using Thermo-Calc software with the TCFE9 database, where the M 23 C 6 and MC carbides dissolved at temperatures of 800°C and 1190°C, respectively. At 900°C, the volume fraction of MC carbide was 0.14. It was clear from the combined experimental findings that the austenitic matrix absorbed more C elements and alloying elements at higher normalizing temperatures. MC carbides were difficult to precipitate with increased austenitizing temperature, according to Ning et al [31]. As a result, during tempering, the MC carbides precipitated, which explained why the volume fraction of MC carbides increased as the normalizing temperature increased. However, the film M-A constituents in the normalized microstructure decomposed into chain-like 10-20 nm carbides, as shown in figure 10(b).
The maximum number density of M 23 C 6 and M 7 C 3 carbides in the tempered microstructure at a normalizing temperature of 900°C was because the undissolved carbides could act as nuclei for precipitation, promoting the coarsening and spheroidization of carbides. As the normalizing temperature increased, fewer nucleation sites remained in the normalized microstructure. The average diameter of large-size carbides grew while their number density declined.

Effect of normalizing temperature on precipitation strengthening
The mechanism of precipitation strengthening in 2.25Cr-1Mo-0.25V steel could be divided into the dispersed strengthening of MC carbides and the precipitation strengthening of large-sized M 23 C 6 and M 7 C 3 carbides. Large-size carbides were frequently located at the prior austenite grain boundaries or bainite lath boundaries for precipitation strengthening. The M 23 C 6 and M 7 C 3 carbides were difficult to bypass and dislocations accumulated around them. The M 23 C 6 and M 7 C 3 carbides efficiently inhibited dislocation movement during tension, which increased the dislocation density within the grains. Precipitation strengthening of the M 23 C 6 and M 7 C 3 carbides could be computed by the dispersed barrier hardening (DBH) model, as follows: where M is the Taylor constant, G is the shear modulus, b is the Burgers vector, N is the number density of largesize carbides, and d 1 is the average diameter of large-size carbides. The values in table 1 for products with an average diameter d and number density N were 366.6, 373.5, 380.3, and 369.6. As the normalized temperature increased, we found that the contribution of large-size carbides toward precipitation strengthening (σ 1 ) varied slightly. However, when the sliding dislocations came into contact with tiny MC carbides, the dislocation line was obstructed and subsequently bent for dispersed strengthening. The dislocations could make their way past the MC carbide barrier as a result of increasing stress, and a dislocation ring formed around the MC carbides. According to the Orowan strengthening mechanism, the dispersed strengthening of MC carbides could be calculated by where f v is the volume fraction of MC carbides and d 2 is the average size of the MC carbides. The volume fraction of MC carbides increased and the average size slightly changed with an increase in normalizing temperature. Consequently, the effect of dispersed strengthening increased, and the dispersion strengthening of MC carbide increased when the normalizing temperature increased from 900°C to 960°C, which was one of the causes of the increase in high-temperature strength.

4.3.
High-temperature fracture mechanism of 2.25Cr-1Mo-0.25V steel The dimples of the sample normalized at 960°C were deeper and larger than those of the sample normalized at 900°C, as indicated by the fracture morphology in figures 11(a) and (b). This demonstrated that an increase in normalizing temperature encouraged the precipitation of MX carbide and enhanced the material's resistance to fracture. Following tensile testing at 550°C, the fracture mechanism was further studied using SEM micrographs of the cross-sectional area of the adjacent fracture. Figures 11(c) and (d) shows that when the experimental steel was under tensile stress, the tensile force lengthened the matrix microstructure in the tensile direction, which led to a larger stress concentration as a result of the high hardness of the carbides. Numerous micro-voids and micro-cracks were observed that started at the large-sized M 23 C 6 carbides and eventually aggregated there. The micro-voids and micro-cracks rapidly spread to surrounding carbide particles during the tensile test, resulting in long cracks that finally caused the material to fracture. Therefore, the large-sized M23C6 carbides were separated from the bainite matrix under tensile stress, which was the cause of high-temperature fracture.

Conclusions
Tempered experimental steels treated at various normalizing temperatures were prepared by replicating the core cooling rate of extra-thick 2.25Cr-1Mo-0.25V steel plates, and their precipitation behavior and mechanical properties were studied. The following results were obtained when the normalizing temperature was raised from 900°C to 990°C.  (1)The microstructure showed the following variations: growth of the original austenite grains, (1) an increase in the original austenite grain size (2) a reduction in the volume fraction of undissolved carbides in the normalized microstructure, with a few high-temperature resistant (Ti,V,Nb)C carbides that were difficult to dissolve into the matrix, (3) an increase in the volume fraction of MC carbides with little change in their average diameter, and (4) an increase in the average diameter of M 23 C 6 and M 7 C 3 carbides with a decrease in their number density.  (2)The high-temperature yield strength increased from 451 to 475 MPa at 550°C. By contrast, the impact absorption energy decreased from 93 to 41 J at −30°C. Thus, 960°C was the ideal normalizing temperature.
(3)Dispersed strengthening played an important role in the precipitation strengthening mechanism. The major cause was that the elevated normalizing temperature encouraged the decomposition of the film M-A constituents and solid solution of alloying elements within the bainite matrix, which in turn promoted the precipitation and formation of MC-type carbides.
(4)The large-sized carbides resulted in a buildup of dislocations, which concentrated stress and caused cracks to develop and expand.
In this work, the experimental results are derived from the experimental samples simulating the cooling rate of the core of the 180mm extra thick 2.25Cr-1Mo-0.25V steel, which has guiding significance for industrial production.

Declarations
Conflict of interest All authors certify that they have no affiliations with or involvement in any organization or entity with any financial interest or non-financial interest in the subject matter or materials discussed in this manuscript.

Ethical compliance
All procedures performed in studies involving human participants were in accordance with the ethical standards of the institutional and/or national research committee and with the 1964 Helsinki Declaration and its later amendments or comparable ethical standards.