Making large-size fail-safe steel by deformation-assisted tempering process | Scientific Reports

Scientific Reports volume 14, Article number: 22345 (2024) Cite this article

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Synergistically improving the strength and toughness of metallic materials is a central focus in the field of physical metallurgy. Yet, there is a noticeable lack of research in strengthening and toughening large-size metal components, whereas those components are extensively used in the modern industry. In this work, a deformation-assisted tempering (DAT) process was proposed to create a novel microstructure in 1.4 tons low-alloyed plain steel. After DAT treatment, the steel contains low dislocation density but high density of low-angle subgrain boundaries and dispersed spherical nano carbides. Such microstructure enables a much better combination of tensile strength and fracture toughness compared to the small-size quench and temper steels. The significant improvement in low-temperature impact toughness is due to the occurrence of delamination and subsequent large plastic deformation at the notch tip. The DAT process can provides a new strategy for the development of large-size fail-safe steel with excellent strength and fracture resistance.

Engineering applications in the fields such as wind energy, marine shipbuilding, and transportation require structural steel that is strong, tough, and reliable, especially when building large-size components1. For instance, the installed capacity of wind turbines used in offshore wind farms has exceeded 10 MW. In automotive industry, the capacity of aluminum alloy die-casting equipment even reaches 10,000 ton. An abundance of sizable, high-quality steel components are required to accommodate these new high-performance equipments2.

Conventional and many new technologies have been developed to improve the strength-toughness balance of steels by microstructure refinement3,4,5,6,7,8. Quenching-tempering (QT) treatment is the most commonly used method to improve the mechanical properties of steels by refining the microstructure through phase transformation9. Besides, reducing the rolling temperature and promoting ferrite nucleation, the so-called Thermo-Mechanical-Control-Process can also refine the microstructure10,11. Furthermore, warm deformation has attracted wide attention as an alternative method for producing high-strength and high-toughness steels with fine-grained microstructure12,13,14. For instance, Kimura et al.15,16 used caliber rolling to deform a medium-carbon low-alloy steel within the tempering temperature range to obtain steel with elongated grain microstructure, which enables an inverse temperature effect on the impact toughness.

However, the aforementioned efforts mainly focused on tailoring the microstructure and mechanical properties of steel products with sizes limited by tens of millimeters, which cannot be used to improve the strength-toughness balance of large-size steel components. Firstly, the low hardenability of low-alloyed steel makes it difficult to obtain a uniformly refined microstructure through phase transformation. In addition, the existing manufacturing technologies are mostly limited to rods and sheets, which are difficult to be used for the preparation of large-size steel.

In this work, a new technology, deformation-assisted tempering (DAT), is proposed. Deformation and long-term tempering under medium temperature conditions result in a novel microstructure consisting of ferrite with nanoscale carbides, high-density low-angle subgrain boundaries (LAGBs), and strong deformation texture. The microstructure and mechanical properties of the 1.4 ton steel sample are investigated in detail. The effects of carbides, grain boundaries, and fracture modes on the strength-toughness balance of the material are discussed.

The 42CrMo continuous casting billet, which was provided by Jiangyin Zhenhong Forging Co., Ltd was selected for the current study and its chemical compositions is listed in Table 1. The steel weights about 1400 kg with a size of 450 × 350 × 1130 mm3 (Fig. 1a). The steel was heated to 1123 K for 12 h, followed by quenching to obtain a fine-grained microstructure. The DAT was carried out in three steps using a bidirectional deformation method: (1) The 42CrMo steel was held at 923 K for 12 h, and then forged (the cross-sectional size of the sample was forged from 450 × 350 mm2 to 250 × 280 mm2). (2) After the single-end forging was completed, the samples were returned to the furnace (923 K) for 12 h, and then forging the clamp end to the same size as the first step (250 × 280 mm2). (3) Returning the sample to the furnace (923 K) for 12 h again, and then forging the sample from 250 × 280 mm2 to 150 × 320 mm2 along the thickness direction. The forged steels were air-cooled to ambient temperature, and the final size of the sample is 150 × 320 × 2100 mm3.

Schematic illustration of forging processes in this work: (a) the size of 42CrMo continuous casting billet; (b) bidirectional deformation process; (c) definition of direction after DAT.

The 4500 ton forging equipment is used. The bidirectional forging process is shown in Fig. 1b. During the whole forging process, the one-pass forging reduction is 70 mm. The total reduction of area after deformation is 67%. To facilitate the discussion below, the direction of the forged sample is shown in Fig. 1c. In addition, the unforged 42CrMo steel cut from the continuous casting billet (100 × 100 × 100 mm3), which was denoted as the QT sample hereafter, was quenched from 1153 K and tempered at 953 K for 210 min to achieve the same tensile strength as DAT steel.

In order to characterize the microstructure of DAT steel, block samples with a cross-section of 10 × 10 mm2 along different plane (Z–X, Z–Y, X–Y) were fabricated from the center and edge of the steel using wire electrical discharge machining (EDM). The microstructures were observed by scanning electron microscopy (SEM, Gemini 300) equipped with an electron backscattered diffractometer (EBSD) detector and transmission electron microscopy (TEM, FEI Talos F200X). The characterization of cementite, high-angle grain boundaries (HAGBs), and LAGBs was carried out using the Transmission Kikuchi Diffraction (TKD) technique performed with an EDAX Velocity Super EBSD detector. The step size used for EBSD and TKD data collection was 0.1 μm and 0.01 μm, respectively. The EBSD and TKD data were analyzed in the AZtecCrystal software. The carbide sizes were analyzed using Nano Measurer software. A non-aqueous electrolyte was employed to extract carbides and maintain their original 3D morphology, more details are available in the literature17.

The uniaxial tensile specimens with a gauge cross-section of 4.0 mm × 2.0 mm and a gauge length of 12 mm were prepared using EDM. Uniaxial tensile tests were performed at room temperature at a crosshead speed of 1 mm/min according to ISO 6892–1: 2019, MOD. The instrumented Charpy pendulum impact tests were conducted using the JBW-450H impact tester in ambient air according to ISO 14556: 2015, MOD. After removing the surface layer of 10 mm using EDM, impact samples were taken at intervals of 10 mm. Four sampling positions were taken from the edge to the center of the plate. The specimen geometrical dimensions were 10 × 10 × 55 mm. At least three valid Charpy impact tests were performed on each microstructural condition to record the dynamic load–displacement curves. The impact tests were conducted at the following temperatures: 77, 113, 153, 193, 213, 233, 253, 273, 293, 313, 353, 393, 433, and 473 K.

The compact-tension specimens were machined in accordance with ASTM standard E182018, with a width of W = 32 mm, thickness of B = 4 mm, and original crack length of am = 12 mm introduced by EDM. Fatigue pre-cracking was performed on all samples prior to the fracture toughness measurement. The crack length of a0 = 16 mm corresponded to the testing requirements of ASTM standard E1820 of 0.45 ≤ a/W ≤ 0.70. The crack extensions during the pre-fatigue process and the fracture toughness test process were monitored by mounting an Epsilon clip gauge (Epsilon technology corp, Jackson, WY, USA) with a gauge length of 5 mm (− 1/ + 4 mm). Fatigue pre-cracks were introduced by tension-tension loading in a 100 kN Servo-Hydraulic fatigue testing machine (LF5105). The cyclic loading was under load control at a stress-intensity range ∆K (Kmax–Kmin) of 20–25 MPa·m1/2 at 10 Hz frequency with a force ratio of R = 0.1. After completing the tests, the length of the fatigue pre-crack region was measured using SEM. The average total crack length was used for subsequent data processing.

The microstructure of DAT steel is shown in Fig. 2. EBSD analysis reveals that the grain morphology is heterogeneous in three dimensions as shown in Fig. 2a. Figure 2b shows representative grain boundary map of DAT sample (X–Z plane), where the blue lines represent HAGBs (θ > 15°) and the red lines represent LAGBs (2° ≤ θ ≤ 15°). The content and density of grain boundaries were quantified based on the method as reported in the literature19, and the results are represented in Table 2. The density of LAGB and HAGB in DAT steel is 868.5 mm−1 and 674.2 mm−1, respectively. The analysis of TKD shows that larger carbides in DAT steel are prone to appear at HAGBs, while smaller carbides are mostly located within the grains and near LAGBs (Fig. 2c). The statistical results show that the sizes of carbides located at HAGBs, LAGBs, and within the grains are 203 ± 86, 120 ± 46, and 62 ± 36 nm, respectively. The Kernel Average Misorientation (KAM) map of the present steel is shown in Fig. 2d and indicates that the regions near LAGBs have much higher KAM values. The DAT process leds to the formation of new fine grains that are free of internal distortions as shown in Fig. S1b in the supplementary material. The grain orientation spread, which is the average deviation in orientation between each point in a grain and the average orientation of the grain20,21, showed that such distortion free grains that can be considered as a results of continuous dynamic recrystallization. These recrystallized grains account for a volume fraction of about 10%, whereas the deformed and substructured grains account for 90%. The pole figure indicates the DAT steel has a strong deformation texture, including Zbcc ({111 < 110 >) and Cube ({001} < 100 >) textures as shown in Fig. 2e. The detailed texture analysis was presented in Fig. S2 in the supplementary material.

Microstructure of the 42CrMo steel manufactured by DAT process: (a) the microstructure reconstructed by EBSD IPF maps; (b) the image quality map superimposed with the grain boundaries. The blue line represents the HAGBs (θ > 15°) and the red line represents the LAGBs (2° ≤ θ ≤ 15°), the insert is the plot of misorientation distribution; (c), (d) TKD analysis results; (e) the pole figures of X–Z plane microstructure.

The carbides of DAT steel are shown in Fig. 3. Figure 3a reveals that the fine carbides in present steel are distributed very uniformly in the ferrite matrix. Note that the prior-austenite grain boundaries are difficult to identify (Fig. 3b). The TEM (Fig. 3c) combined with SEM (Fig. 3d) observations reveal that most of the carbides are spherical. The selected area electron diffraction (SAED) pattern of the carbide corresponds to Fe3C (Fig. 3e). Figure 3f. is a high-resolution image of the interface between spherical carbide and matrix. The interfacial strain was derived from the high-resolution image by geometric phase analysis (GPA) method22. The strain component εxx and εxy are shown in Fig. 3g,h, respectively. The localized strain concentration that appeared around the interface suggests the non-coherence interphase boundary.

(a), (b) The microstructure reconstructed by SEM maps; (c), (e) bright-field TEM image of the present steel, the insert is a SAED pattern of the carbide corresponding to Fe3C; (d) the morphologies of the carbides as observed by SEM; (f) high-resolution TEM image showing the interface between carbide and ferrite matrix; (g), (h) the strain map corresponding to (f) derived by GPA software.

For comparison, the microstructure of QT steel is shown in Fig. 4. The microstructure of 42CrMo after QT treatment exhibits tempered martensite, without any discernible texture as shown in Fig. 4a. The QT structure contains numerous HAGBs (Fig. 4b), with carbides primarily precipitating along the prior-austenite grain boundaries (PAGBs) and the lath martensite interface (Fig. 4c,d). Carbides exhibit a morphology of both granular and elongated shapes (Fig. 4e). Figure 4f illustrates that a substantial amount of Cr and Mn elements in the carbide.

Microstructure of QT 42CrMo steel: (a) IPF map; (b) grain boundary map; (c) SEM image; (d) bright-field TEM image; (e) the 3D morphologies of the carbides obtained by electrolytic extraction; (f) elemental mapping results show that the cementite phase in QT sample is enriched in Cr and Mn.

Figure 5a shows the engineering tensile stress–strain curves of DAT steel and QT steel. The tensile properties are summarized in Table 3. The yield strengths (σy) of the specimens prepared along the X, Y, and Z direction are 640 MPa, 560 MPa, and 512 MPa, respectively. The ultimate tensile strength (σB) of DAT steel is almost equivalent to that of QT steel, but the uniform elongation (εu) and local elongation (total elongation (εt)—εu) of DAT steel are much larger than those of QT steel. It has been reported that the larger ductility in ferritic steel can be attributed to the improved strain hardening rate as a result of the fine carbides23.

The fracture toughness of present steels is determined by nonlinear elastic fracture mechanics methods as shown in Fig. 5b. The intersection of the resistance curve with the 0.2 mm offset/blunting line with a slope of 2σ0 (σ0 = 0.5(σy + σB)) defines a provisional toughness JQ. The results show that the JQ of QT steel is 303 kJ/m2, while the JQ of DAT steel along X, Y, and Z directions are 932 kJ/m2, 681 kJ/m2, and 547 kJ/m2, respectively. Therefore, it can be noted that the fracture toughness of DAT steel in three directions is superior to that of QT steel under the same strength.

(a) Engineering stress–strain curves of the present steel deformed under tension along the X, Y, and Z directions; (b) the J-integral-based resistance curves (J-R curves) showing the variation in the J-integral as a function of crack extension Δa for different specimens.

Figure 6 shows the results of the Charpy V-notch impact properties (vE) vs temperature for the DAT and QT samples. Two types of impact specimens were processed (Type-1: V-notch direction perpendicular to Z; Type-2: V-notch direction parallel to Z). Additionally, the definitions of the crack types observed after fracture in these two types of impact specimens are illustrated in Fig. S3 of the supplementary material24,25,26,27. The ductile–brittle transition temperature (DBTT), defined as the temperature corresponding to half of the upper-shelf energy, of the type-1, type-2, and QT samples are 180 K, 223 K, and 263 K respectively. The DAT samples have a higher upper-shelf energy than QT samples. Furthermore, the type-1 impact specimen exhibited an inverse temperature dependence of toughness within the range of 193–293 K. In other words, the impact toughness of DAT steel increases with decreasing temperature at temperatures ranging from 193 to 293 K. In the upper-shelf region of the steel, all three types of impact samples exhibit ductile fracture, with cracks propagated directly across the central portions of the impact specimens. As the testing temperature decreases, delamination cracks perpendicular to the impact load appear in the type-1 specimen, and divider cracks gradually increase in the type-2 specimen, while the brittle region in the QT steel gradually expands. The fracture morphology and the corresponding load displacement curves can be observed in Figs. S4 and S5 of the supplementary materials.

Changes in the Charpy V-notch impact energy as a function of testing temperature for different steels.

Figure 7 presents the crack-path profile of DAT steel at 293 K, 233 K, and 193 K. At room temperature (Fig. 7a), the cracks are parallel to the load direction. At 233 K (Fig. 7c), the multiple delamination cracks can be observed. Those are perpendicular to the load direction and connected by vertical cracks (approximately 45° to the load direction). At 193 K (Fig. 7e), the delamination cracks only appear near the notch. After delamination, the sample undergoes significant plastic deformation without being broken. It has been reported that cracks delaminate at weak interfaces, such as inclusions, phase interfaces, interfaces between cementite and ferrite, and grain boundaries24,27,28. Furthermore, delamination can also occur along cleavage surfaces. The morphology of the delamination layer in Fig. 7f is characterized by a river-like cleavage fracture pattern.

Crack-path profile captured on the mid-plane section of the DAT specimens (Type-1) at different impact temperatures: (a), (b) ambient temperature; (c), (d) 233 K; (e), (f) 193 K.

After DAT process, the microstructure (Figs. 2 and 3) of large-sized 42CrMo steel contains dispersed nano carbides, high density of LAGBs, and exhibits strong deformation texture. During the DAT process, accumulated plastic strain and prolonged tempering are key factors for the formation of such microstructure. Firstly, the initial dominant microstructure of 42CrMo steel (after heat treatment for forging) is ferrite plus pearlite since the large piece cannot be quenched to martensite (the continuous cooling transformation curves and corresponding microstructures of 42CrMo steel are given in Fig. S6 and S7, respectively). The cementite lamellae in pearlite disintegrated into fine fragments during the warm forging, and these fine fragments can be spheroidized into discrete cementite particles29. The driving force for spheroidization of carbide depends on the reduction of interface area between ferrite and cementite lamellae30,31.

The redistribution of nanoscale carbides during DAT is related to the diffusion of dissolved carbon along the dislocations and grain boundaries29,32. During the warm forging process, a large number of dislocations are stored due to the accumulation of strain that can provide a pathway for dissolved carbon to diffuse from the former pearlite colonies to ferrite. During the subsequent long-term tempering process, the carbides precipitate from the ferrite while the dislocations annihilate30. Therefore, the distribution of carbides in both the center and edge of the DAT sample is relatively uniform (Fig. 3b and S1(d)). A relatively uniform microstructure results in minimal differences in mechanical properties (Fig. S8). In contrast, the carbides in 42CrMo steel after QT treatment are precipitated along the PAGBs and lath martensite interface (Fig. 4c, d). Most carbides exhibit a striped shape as illustrated in Fig. 4e.

The formation of texture is related to the history of forging. There are two types of textures in DAT steel, namely Zbcc texture and Cube texture. The strength of Zbcc texture is greater than that of Cube texture. Warm forging adopts a bidirectional forging method, which first forms a cube texture, similar to the texture of the sample prepared by caliber rolling24,33,34. The final pass of warm forging forms Zbcc texture. Furthermore, the formation of LAGBs is related to the evolution of dislocation substructure. During the DAT process, the accumulation and recovery of dislocations forms dense dislocation walls (DDWs). Interaction of more dislocations with DDW results in progressive sharpening of DDW and then gradually evolve into LAGBs35. The fractions of HAGBs and LAGBs in DAT sample are almost the same, while the QT microstructure is mainly composed of HAGBs.

The DAT steel has high quasi-static fracture toughness and impact toughness at room temperature. The fracture toughness of the DAT sample in all three directions is higher than that of the QT sample as shown in Fig. 5b. It indicates that the DAT sample is less prone to fracture under static and quasi-static conditions. Furthermore, the room temperature impact toughness of DAT samples is higher than that of QT samples. The room temperature impact toughness of DAT steel is higher than 150 J as shown in Fig. 6. Figure 8 schematically illustrates the difference in intrinsic toughening mechanism between QT steel and DAT steel. The QT sample contains higher density of HAGBs, that is as large as 2274.8 mm−1. The dislocations can be quickly piled up at these HAGBs, resulting in high stress concentration. Furthermore, the necklace-like carbides at HAGBs are prone to reducing the cohesion of grain boundaries, thereby facilitating crack initiation at these locations and leading to the occurrence of cleavage fracture36,37,38,39,40 (Fig. 8a). Numerous carbides have been observed at the PAGBs and lath martensite interfaces in the fracture surface of the QT steel as shown in Fig. S9 in the supplementary material. Therefore, the QT steel has relatively low toughness. For comparison, the DAT steel has much lower density of HAGBs but higher density of LAGBs. The LAGBs provide a little hinderance effect against the movement of dislocations. It has been reported that the interactions between LAGBs and dislocation during plastic deformation can reduce the internal stress, and improve the fracture toughness of the steel41,42. Besides, the fracture toughness and resistance to delayed fracture can be improved by the presence of fine spherical carbides43,44,45,46,47,48. During plastic deformation, more slip systems at the interface between spherical carbides and the ferrite matrix can be triggered. So that the synergistic plastic deformation ability between spherical carbides and the matrix is higher (Fig. 8b) in the DAT steel as compared to the QT steel.

Schematic illustration showing intrinsic toughening mechanisms associated with grain boundary, carbide, and texture: (a) QT steel; (b) DAT steel.

Furthermore, the DAT samples have strong {111} < 110 > deformation textures. Numerous studies have shown that the extraordinary toughness of warm deformed steel is related to the presence of < 100 > crosswise RD texture formed during the deformation process24,25,49,50. Xie et al.51 used molecular dynamics simulations investigate the toughness mechanisms of the ultrafine grain structure steel with different textures. The results show that the sliding of grain boundaries({001}/{110} type and {110}/{111} type) can arrest the cracks which propagate along the load direction and improve toughness. Therefore, the intrinsic toughening mechanisms through closely spaced LAGBs, dispersed nano carbides, and strong texture can significantly increase both the crack-initiation and crack-growth toughness of DAT steel52.

Delamination is another important factor that improves the toughness of DAT steel. According to the Yoffee diagram, brittle fracture occurs when the peak tensile stress (σt) in the process zone of a crack tip exceeds the brittle fracture stress (σf)24,50. In BCC materials, σy (which is proportional to σt) significantly increases with decreasing temperature25,53. Due to the anisotropy of grain morphology, the σf⊥X become smaller than σf//X54,55. Additionally, DAT steel exhibits strong deformation texture, providing numerous {100} cleavage planes in the X direction. Therefore, as the temperature decreases, the delamination phenomenon in DAT steel becomes more pronounced.

To investigate the delamination effect on the DBTT, the impact load and deflection during the impact test were recorded continuously as a function of time using an instrumented Charpy impact test as shown in Figs. S4 and S5 in the supplementary materials. Figure 9 shows the energies consumed during different stages of fracture for type-1, type-2, and QT samples at different impact temperatures. Figure 9a gives a graphical illustration of the absorbed impact energy values corresponding with crack growth and fracture. Figure 9b shows a schematic diagram of the compliance changing rate (CCR) method56,57. The details of the analysis method are described in the supplementary material.

Figure 9c–e and Table 4 presents the impact absorption energies corresponding to different fracture stages of the three types of impact specimens. It is evident that during the crack initiation stage (Win), type-1 and type-2 samples exhibit higher absorption energies compared to QT samples, indicating that DAT steels possess greater resistance to crack initiation. For type-2 and QT samples, the resistance to crack growth (Wstable + Wunstable) decreases with decreasing the testing temperature. Once cracks initiate, the samples immediately fracture when test temperature is below the DBTT. Wstable of the type-1 sample decreases with a decrease in testing temperature, but it remains significantly higher than that of the QT sample at the same temperature. It is intriguing to note that the Wunstable of type-1 samples increases with decreasing test temperature, reaching a maximum of 337 J at 193 K.

(a) Schematic diagram of absorbed energies at different fracture stages; (b) illustration of the CCR method for determining the load at crack initiation Pin; (c), (d), (e) absorbed impact energies associated with different stages for type-1, type-2, and QT specimens subjected to instrumented Charpy impact testing with different impact temperature.

The DAT steel (Type-1) exhibits a significant inverse temperature dependence of toughness in the range of 193–233 K. Within this temperature range, the impact specimens experienced crack delamination and did not completely fracture. There are two factors that control the low-temperature toughness of DAT steel: (1) The delamination process generates two new interfaces that release a certain amount of strain energy. The occurrence of delamination is believed to alleviate the high triaxial stress conditions ahead of the advancing crack front, thereby resulting in high impact energy. (2) The unbroken part underwent a pure bending process, which consumes a significant amount of energy by plastic deformation. The plastic deformation ability of DAT steel at low temperatures is related to the nucleation and propagation of dislocations. Compared with QT steel, the DAT steel has a lower initial dislocation density and lower resistance to the movement of new dislocations. The large number of grain boundaries and subgrain boundaries in DAT steel provide sites for the nucleation of dislocations58,59.

In this work, a large-size 42CrMo steel with outstanding strength-toghness synergy capability was developed. The deformed microstructure and toughing mechanisms were systematically investigated and discussed. The main findings can be summarized as:

(1)A heterogeneous microstructure composed of dispersed nano carbides, large number density of LAGBs, and strong deformation texture was formed in 1.4 tons large-sized 42CrMo steel using DAT technology.

(2)The present DAT large-size 42CrMo steel demonstrates an excellent combination of strength, ductility, and toughness. The strength-ductility balance (σB × εt) of the DAT steel is 22,194 MPa·%, while that of the QT steel is 13,405 MPa·%. The fracture toughness (JQ) for crack initiation can be as high as 932 kJ/m2, while the fracture toughness of the QT steel is 303 kJ/m2. The room temperature impact toughness of the DAT steel exceeds 150 J. Furthermore, the impact toughness of DAT steel at 193 K is as high as 300 J, while the impact toughness QT steel at this temperature is lower than 30 J.

(3)The extremely high toughness of the present DAT steel can be attributed to both the intrinsic toughening and extrinsic toughening mechanisms. The intrinsic toughening mechanism involves the high plasticity capacity of the DAT steel due to the low initial dislocation density, dispersed nano carbides, and strong texture. The extrinsic toughening mechanism refers to the presence of delamination, which effectively releases a large amount of elastic energy and reduces the locus stress concentration at the crack tip.

The data that support the findings of this study are available on request from the corresponding author, Zhichao Luo，upon reasonable request.

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The authors acknowledge the financially supported by National Key R&D program [No. 2022YFB3707501], GDAS' Project of Science and Technology [2022GDASZH-2022010202], Guangdong Provincial Project [2022A0505050053]

Research Institute for Energy Equipment Materials, Tianjin Key Laboratory of Materials Laminating Fabrication and Interfacial Controlling Technology, School of Materials Science and Engineering, Hebei University of Technology, Tianjin, 300132, China

Kuanyuan Fan, Baoxi Liu & Fuxing Yin

Institute of New Materials, Guangdong Academy of Science, Guangdong Provincial Key Laboratory of Metal Toughening Technology and Application, Guangzhou, 510651, China

Kuanyuan Fan, Tianlong Liu, Fuxing Yin & Zhichao Luo

Belgorod State University, Belgorod, Russia, 308015

Andrey Belyakov

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K.F. and T.L. carried out the experiment. K.F. wrote the manuscript with support from Z.L. and B.L. F.Y. and A.B. helped supervise the project and contributed to the interpretation of the results. Z.L. took the lead in writing the manuscript. All authors provided critical feedback and helped shape the research, analysis and manuscript.

Correspondence to Baoxi Liu or Zhichao Luo.

The authors declare no competing interests.

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Fan, K., Liu, B., Liu, T. et al. Making large-size fail-safe steel by deformation-assisted tempering process. Sci Rep 14, 22345 (2024). https://doi.org/10.1038/s41598-024-70576-3

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Received: 01 June 2024

Accepted: 19 August 2024

Published: 27 September 2024

DOI: https://doi.org/10.1038/s41598-024-70576-3

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