ISIJ International, Vol. 61 (2021), No. 1

© 2021 ISIJ463

* Corresponding author: E-mail: denglei@hust.edu.cn

© 2021 The Iron and Steel Institute of Japan. This is an open access article under the terms of the Creative Commons Attribution-NonCommercial-NoDerivs license (https://creativecommons.org/licenses/by-nc-nd/4.0/).

ISIJ International, Vol. 61 (2021), No. 1, pp. 463–472

https://doi.org/10.2355/isijinternational.ISIJINT-2020-397

1. Introduction

Ultrahigh strength steels are widely used in heavy load-bearing components, such as aircraft landing gears and drive shafts owing to its ultra-high-strength and resistance to stress corrosion.1) Due to the relatively low cost, simple production process and good comprehensive performance, medium-carbon and low-alloy ultrahigh strength steels are the most widely used among the ultrahigh strength steels.2,3) Besides higher strength, better toughness and longer fatigue life have been always the urgent demands of these steels, especially for components used in high-impact and long-term service environments. Generally, the increase of retained austenite (RA) can improve the toughness of steels, but it would result in a decrease of strength that is considered to have a positive correlation with fatigue life. Increasing the content of bainite is another effective method to improve the toughness of steels, similarly, it also comes in the expense of strength. From these results, a common contradictory phenomenon can be recognized that the

Simultaneous Improvement of Toughness and Fatigue Life in a Typical Ultrahigh Strength Steel by a New Deep Cryogenic Treatment Process

Peng GUO, Lei DENG,* Xinyun WANG and Junsong JIN

State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, 1037 Luoyu Road, Wuhan, Hubei, 430074 China.

(Received on July 8, 2020; accepted on September 3, 2020)

As a supplemental process of heat treatment, deep cryogenic treatment (DCT) is believed to have favor-able effects on the mechanical properties of steels. In this study, three different sequences of DCT and tempering, including QCTT (Q: quenching, C: DCT, T: tempering), QTC, and QTCT, were investigated to discover an effective method to improve the impact toughness and fatigue life of a typical ultrahigh strength steel simultaneously. Compared with the traditional heat treatment process, the impact tough-ness, fracture elongation, and fatigue life were improved by 10%, 17.8%, and 13.7%, respectively, after the QTC treatment process. The microstructures were characterized by scanning electron microscope (SEM), electron backscatter diffraction (EBSD), transmission electron microscopy (TEM), and electron probe micro-analysis (EPMA). The results indicate that after QTC process, the formation of cracks was restricted due to the decrease of carbides and disappear of twinned martensite, both of which would serve as the sources of cracks. Besides, the propagation of cracks is hindered because the decrease of blocky retained austenite and increase of filmy retained austenite. As a result, the impact toughness, fracture elongation and fatigue life are improved simultaneously. This work provides a new and effective method to simultaneously improve the toughness and fatigue life of ultrahigh strength steels.

KEY WORDS: deep cryogenic treatment; toughness; fatigue life; retained austenite; carbides.

improvement of toughness is typically achieved with the sacrifice of strength.4)

During the past few decades, many studies have been carried out to improve the toughness and fatigue life of ultrahigh strength steels. By increasing the content of bainite via the austempering process, Tomita et al. improved the fracture toughness of 300M steel but decreased the mate-rial strength.5) Another method of increasing the austen-itizing temperature employed by Ritchie et al. improved the fracture toughness of AISI 4340 without reducing the material strength but with an apparent decrease in impact energy.6) Additionally, a rapid cyclic austenitizing method of reducing the prior austenite grain size proposed by Sastry et al. and Padmanabhan et al. improved the strength and impact toughness of a low-alloy ultrahigh strength steel.7,8) However, the use of rapid austenitization is difficult for large components due to the longtime heat transfer from surface to interior. To improve fatigue life, Amrita et al.9,10) employed a method of shot peening that increased the residual compressive stress, but this method becomes ineffective when the applied stress exceeds 0.64 times of yield stress. At present, there still has been no relatively

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satisfactory way established to simultaneously improve the toughness and fatigue life of large ultrahigh strength steel components. Thus, it is necessary to discover new methods for improving the toughness and fatigue life.

Deep cryogenic treatment (DCT), which has been intro-duced in the heat treatment process of some steels, has been verified as capable of improving the toughness and fatigue life of steels. It is performed after quenching and prior to tempering which can increase the undercooling of austenite. As a supplemental process of heat treatment, DCT can promote the transformation of retained austenite (RA) into martensite and optimize the distribution and morphol-ogy of carbides.11–13) Thus, DCT is always used for some high-carbon or high-alloy steels, and obtained an increase in strength, hardness, or wear resistance.14–19) Different from these steels, the ultrahigh strength steels, such as 300M steel, have less RAs at room temperature duo to the higher Ms. To obtain a relative good comprehensive performance, the conventional heat treatment process of medium-carbon and low-alloy ultrahigh strength steels is oil quenching followed by multiple tempering processes,20) which is a compromise of toughness and strength. This process is also different from the conventional quenching-tempering process mentioned above in conventional steels, and the influence of the introduction of DCT on the mechanical properties is still unknown. It may be simply speculated from the existing researches that the mechanical properties will not be improved significantly when treated by the con-ventional DCT method because of the relatively low volume fraction of RA.21) Recently, study of adjusting the sequence of DCT during heat treatment has found further improve-ment of mechanical properties of a carburized steel.22) Thus, adjusting the sequence of DCT during heat treatment may be an effective way to improve the mechanical properties, which also needs to be investigated.

In present study, three heat treatment processes with different sequences of DCT and tempering were investi-gated. The effects of DCT on microstructures evolution and mechanical properties of a typical ultrahigh strength steel were discussed.

2. Experimental Procedure

The chemical composition (wt%) of the as-received medium-carbon and low-alloy ultrahigh strength steel mea-sured by an X-Ray fluorescence spectrometer (Zetium) and a carbon-sulfur analyzer (CS-3000G), as 0.39C-0.81Mn-1.68Si-1.87Ni-0.85Cr-0.08V-0.42Mo-(Bal)Fe. Compared to AISI 4340, the increase of Si can hinder the transformation of carbide into cementite, which can improve the tempering temperature up to 315°C without embrittlement.23) However, when the tempering temperature is higher than 310°C, the toughness does not improve obviously, but the strength decreases significantly.3) The recommended tempering temperature of this ultrahigh strength steel is 302 ± 6°C,24) which can obtain an ideal state of phase distribution with the best performance. In this study, the tempering temperature is set as 300°C. The initial microstructure is a typical com-pound of ferrite and pearlite after annealing. A traditional heat treatment process and three different sequences of DCT and tempering treatment were performed, as shown

in Fig. 1. First, all samples were austenitized for 1 h at 870°C and oil quenched. The quenched samples were then tempered twice at 300°C for 2 h, followed by air cooling after each tempering process. This process is denoted as quenching-tempering-tempering (QTT) with the solid line shown in Fig. 1(a). The dash line in Fig. 1(a) represents the conventional DCT process of quenching-cryogenic-tem-pering-tempering (QCTT). In this treatment process, after oil quenching, the samples were slowly cooled to −196°C (−3.3°C min −1) for DCT and remained for 12 h. The sam-ples were then heated to room temperature (25°C) at a rate of 3.3°C min −1. Finally, the samples were tempered twice, similar to QTT. The dash dot line in Fig. 1(b) represents the new DCT process of quenching-tempering-cryogenic (QTC), in which the quenched samples were first tempered and then subjected to DCT. The last process is an exten-sion of QTC, with a second tempering process, denoted by quenching-tempering-cryogenic-tempering (QTCT), as shown in Fig. 1(b) with the short dot line.

After different treatment processes, the Rockwell hard-ness of the treated samples was obtained using a Rockwell hardness tester (600 MRD). Uniaxial tensile tests were carried out on a mechanical testing machine (AG-IC 100 kN) according to GB/T228.1 to obtain tensile strength

Fig. 1. Different treatment processes: (a) quenching-tempering-tempering (QTT) and quenching-cryogenic-tempering-tempering (QCTT); (b) quenching-tempering-cryogenic (QTC) and quenching-tempering-cryogenic-tempering (QTCT). (Online version in color.)

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and fracture elongation. V-notch impact tests at 25°C and −40°C were carried out on a pendulum impact tester (SANS ZBC4000) according to GB/T229 to obtain the impact energy. Fatigue lives at various stresses were obtained using a fatigue testing machine (EHF-UV100k2-040-1A) accord-ing to GB/T3075.

After mechanically grinding and polishing the sample surfaces, the micromorphology was determined using a solution of 4% nitric acid and using scanning electron

microscopy (SEM, Quanta 650 FEG) with an electron backscatter diffraction (EBSD) probe. Phase content and RA morphology were obtained through EBSD tests with a scan step of 0.1 μm. The EBSD samples were electro-polished in an A2 electrolyte at 20 V for 15 s. The distribution and morphology of carbides were observed with electron

Fig. 2. The mechanical properties of samples treated by different processes: (a) ultimate tensile strength and Rockwell hard-ness; (b) V-notch impact energy and fracture elongation; (c) fatigue life-stress curves of samples treated by QTT and QTC. (Online version in color.)

Fig. 3. SEM micrographs of samples treated by (a) QTT, (b) QCTT, (c) QTC, (d) QTCT. M: martensite; B: bainite, RA: retained austenite. (Online version in color.)

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probe micro-analysis (EPMA-8050G) in the backscattered electron (BSE) imaging mode, and elemental analysis was determined using the attached wavelength dispersive X-ray (WDX) spectroscopy, EPMA-8050G. Detailed morpholo-gies of martensite and RA were observed using field emis-sion transmission electron microscopy (FETEM, FEI Tecnai G2 F20). The TEM samples were prepared by twin-jet electro-polishing at −30°C and 15 V in a solution of ethanol compound with 8% perchloric acid.

3. Results and Discussion

3.1. Mechanical PropertiesThe mechanical properties of samples treated by differ-

ent processes are shown in Fig. 2. Compared with QTT, the strength and hardness of the QCTT samples are slightly decreased, with a slight increase in the samples treated by QTC and QTCT. Impact energy and fracture elongation shown in Fig. 2(b) exhibit the same trend. The QCTT sam-ples have the lowest impact energy (14.3 J at 25°C and 12 J at −40°C) and fracture elongation (7.85%), while the QTC samples have the highest values (23.7 J, 17.7 J, and 9.27%, respectively). The performance of the QTCT samples (22.3 J, 16.7 J, and 8.77%, respectively) are between the QTC and QTT samples (21.7 J, 16 J, and 7.87%, respectively). The

fatigue life of the QTC samples is significantly higher than the QTT samples, as shown in Fig. 2(c). These results indi-cate that DCT before tempering has adverse effects on the strength and toughness. However, when DCT is introduced after tempering or between two tempering processes, the impact toughness and fracture elongation are significantly improved, with only a slight increase in strength and hard-ness. Furthermore, the fatigue life of the QTC samples is also improved greatly compared with the QTT samples.

3.2. Microstructures and Fracture MorphologiesThe SEM micrographs of different treatment processes

are shown in Fig. 3. The main phase is martensite, with a small amount of RA, bainite, and carbides. In the QTT sample shown in Fig. 3(a), many blocky RAs (presents a uniform black color, as shown in the high magnification graph in Fig. 3(a)) and a minor amount of bainite (in the form of a double lens, and the inside is short rod-like struc-ture composed of ferrite and carbide, as shown in the high magnification graph in Fig. 3(b)) can be observed. When introducing DCT after quenching and prior to tempering (QCTT), the content of bainite has increased, and parts of the blocky RAs have decomposed, as shown in Fig. 3(b). In addition, when introducing DCT after quenching and tem-pering (QTC), the amounts of carbides and bainite are also

Fig. 4. SEM micrographs of impact fracture morphologies at 25°C for (a) QTT sample, (b) QCTT sample, (c) QTC sample; at − 40°C for (d) QTT sample, (e) QCTT sample, (f) QTC sample; and (g) and (h) are the micrographs of Zone I and Zone II in (e), respectively. (Online version in color.)

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reduced, as shown in Fig. 3(c). Meanwhile, minimal blocky RAs can be observed, which have been replaced by filmy RAs. As shown in Fig. 3(d), with the additional tempering process after QTC, the number of carbides has increased, indicating that the tempering process can promote the pre-cipitation and aggregation of carbides.

The SEM micrographs of the V-notch impact fracture morphologies are shown in Fig. 4, in which (a), (b), and (c) had been tested at 25°C for the QTT, QCTT, and QTC samples, respectively; and (d), (e), and (f) had been tested at −40°C. All the fracture surfaces consist of notch, shear lip, and radial region. However, when comparing the fracture surfaces at 25°C, the QCTT sample has the smallest shear lip area, with the lowest corresponding impact energy. This phenomenon is consistent with the theory that the shear lip ratio has a positive correlation with impact energy.25) Differ-ent from 25°C, two different zones in the radial region are observed in the fracture surfaces at −40°C. One is a bright zone that consists of numerous typical dimples, indicating that the fracture mode of region I is a ductile fracture, as shown in Fig. 4(g). The other is a gray zone, with some quasi-cleavage fracture planes in addition to the dimples, as shown in Fig. 4(h). These quasi-cleavage fracture character-istics demonstrate that local brittle fractures appear at low temperatures. The bright area of the QCTT sample is signifi-cantly larger than in the QTC sample when comparing Figs. 4(d)–4(f), and the corresponding impact energy of QCTT is significantly lower than that of QTC. These results suggest that the introduction of DCT after quenching and tempering can reduce the brittle fracture ratio at low temperatures and improve the impact toughness.

The SEM micrographs of tensile fracture morphologies are shown in Fig. 5, presenting the fracture surfaces of samples treated by QTT (Fig. 5(a)), QCTT (Fig. 5(b)), QTC (Fig. 5(c)), and QTCT (Fig. 5(d)). All the fracture surfaces treated by different treatment processes present ductile fracture characteristics, with slight differences among them. More carbides are observed in the fracture surfaces of the QTT and QCTT samples than in the other samples. Additionally, dimple characteristics, evident quasi-cleavage characteristics, micro-cracks, and microvoids are also observed, as shown in Figs. 5(a) and 5(b); the amount of carbides has significantly decreased in the QTC and QTCT samples. Furthermore, negligible large-area quasi-cleavage features appear, and fracture surfaces are present in the form of dimples. It can be summarized that two consecutive tem-pering processes can significantly promote the precipitation of carbides, which serve as sources of micro-cracks in the initial fracture stage.26,27) Quasi-cleavage planes are likely to develop during the fracture process when bands of large-sized carbides have formed,28–30) leading to a significant drop in toughness and elongation.

3.3. Discussion for Mechanism of Toughness and Fatigue Life Optimization

For the fracture behaviors of impact, tensile, or fatigue, the essence is the formation and propagation of cracks in the material. The primary factors affecting the fracture of carbon steel are both the content and morphology of mar-tensite, RA, and carbides.31–34) Thus, the causes of the varied mechanical properties from the different treatment processes

were analyzed through the microstructures. In the follow-ing, the process of traditional heat treatment (QTT), the process with the worst mechanical properties (QCTT), and the process with the best mechanical properties (QTC) will be discussed in detail. The process of QTCT which presents

Fig. 5. Tensile fracture morphologies of samples treated by (a) QTT, (b) QCTT, (c) QTC, and (d) QTCT. (Online version in color.)

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the medium mechanical properties is an extension of QTC, and the main difference is the increase of carbides compared with the QTC due to the additional tempering process, as seen in Figs. 3 and 5. The increase in carbides can lead to an decrease in toughness.35,36) Thus, the QTCT process will be discussed curtly.

Despite this steel is a medium-carbon and low-alloy steel, austenite cannot transform into martensite completely after quenching, and it is distributed among martensite in the form of RA whose content and morphology have important influences on the mechanical properties.37) DCT can pro-mote the transformation of RA into other phases, leading to a decrease in RA content. As a result, the highest RA con-tent appears in the QTT sample (4.7%), and the lowest is in the QCTT sample (1.34%), as shown in Fig. 6. These results can be explained by the equation of fA = exp[−1.10 × 10 −2(Ms−Tq) ],38) which shows the relationship between the RA volume fraction and temperature. In the equation, fA is the volume fraction of retained austenite, and Tq is the low-est tempering reached. It can be deduced that with decreas-ing temperature caused by DCT, the fraction of austenite decrease, indicating that RA continuously transforms at an ultra-low temperature. Generally, the decrease in RA will result in an increase in material strength.39,40) However, the QCTT sample with the lowest RA content has the lowest strength, as seen in Fig. 2(a). The reason for this abnormal phenomenon may be the transformation of bainite, which is a softer phase compared with martensite. When DCT is car-ried out after quenching and prior to tempering, unconverted RA will transform into bainite in the subsequent tempering process at the bainite transformation temperature range due to the segregation of the C atoms.41) As a result, the mar-tensite does not increase correspondingly although the RA content is reduced, and a considerable part of RA transforms into bainite, leading to the decrease in strength. However, when quenching is followed by tempering first, the behavior of RA and C become different. This process is similar to the Q&P process.42,43) During partitioning process, C atoms diffuse from martensite to austenite, increasing the carbon content of austenite. As a result, the stability of austenite is increased, and austenite can exist stably even at room tem-perature. Similarly, during tempering process before DCT, C atoms also diffuse from martensite to RA, increasing the carbon content of RA, which can improve the stability of RA. In the subsequent DCT process, some RA continuously

transforms into martensite due to greater undercooling, and the stability of unconverted RA is improved further. Differ-ent from the Q&P process, the content of RA in this study is less, as a result, C atoms precipitated from martensite during tempering process cannot be completely absorbed by the surrounding RA, and part of the C atoms combine with alloy elements to form carbides. In addition, a considerable part of C atoms remain in martensite to ensure the strength of martensite. In this way, the strength is maintained and the toughness is improved.

However, the decrease in RA will reduce the impact energy, as can be seen from the comparison of the QTT and QCTT samples in Fig. 2(b). In general, blocky RA and filmy RA are the common forms of RA in the martensitic matrix, and they influence fractures differently.44,45) Blocky RA is readily transformed into strain-induced martensite due to the poor mechanical stability,46) aggravating the stress concentration. Secondary cracks are easily generated, which accelerates the crack propagation.47) Slight plastic deforma-tion can effectively relieve the stress concentration at the crack tip because of the good plasticity and mechanical sta-bility of filmy RA, which partially hinders the propagation of cracks. In the QTT sample, RA exists mostly in the form of blocky RA. While in the QTC sample, the content of RA is 3.08%, with the blocky RAs mostly disappearing, and the fraction of filmy RA increases, as seen in Fig. 6(c). A detailed morphology of blocky and filmy RAs is presented in the TEM images in Fig. 7. The dark-field image in Fig. 7(a) illustrates that in the QTT sample, numerous blocky RAs with various sizes are distributed among martensite, and the selected area electron diffraction (SAED) pattern (as shown in Fig. 7(a)) verifies the existence of austenite. In the QTC sample, filmy RAs with different widths are distributed among the martensite, as can be seen in Fig. 7(b). These filmy RAs are formed from blocky RAs due to the influence of DCT after tempering, with all these formations positively influencing the improvement in toughness. Particularly, twinned martensite (TM) is observed in the QCTT sample, as shown in Figs. 7(c) and 7(d), and the SAED pattern shown in Fig. 7(d) verifies the existence of TM. The driv-ing force of C atoms precipitating from martensite increases due to volume shrinkage and increasing C oversaturation caused by the ultra-low temperatures when DCT is intro-duced before tempering. However, because of the difficulty in the diffusion of C atoms at ultra-low temperatures and

Fig. 6. EBSD maps of phase distribution and content: (a) QTT, (b) QCTT, and (c) QTC. M: martensite; RA: retained austenite. (Online version in color.)

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Fig. 7. TEM images of (a) dark-field image and selected area electron diffraction (SAED) pattern in the QTT sample, γ represents FCC type structure, (b) filmy RA in the QTC sample, (c) TM in the QCTT sample, and (d) morphol-ogy and SAED pattern of TM, α represents BCC type structure. (Online version in color.)

Fig. 8. EPMA results of the QTT sample show (a) the BSE image and elemental distribution of (b) C, (c) Mo, (d) Cr, (e) Ni, (f) Si, (g) V, (h) Mn, and (i) Fe. (Online version in color.)

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the short diffusion distance, C atoms would accumulate in the surrounding RAs.48) As a result, the RAs with increased C content transform into TMs in an ultra-low temperature environment.49) The toughness of TM is very poor due to its high C content.50) Cracks are easily formed in TM because of the stress concentration when external loads are applied. It can be concluded that the TM generated in the QCTT process is also the reason for decreasing the impact energy in addition to the decrease in RA content.

Carbide is another significant factor affecting mechanical properties. It can be found that the number of carbides in the QTT and QCTT samples is substantially higher than in the QTC and QTCT samples, and the corresponding impact energy and fracture elongation are lower due to the carbides serving as micro-crack sources. The carbide compositions could be obtained by the EPMA results, in which the mor-phology and elemental distribution can be clearly seen from the BSE image (Fig. 8(a)) and spectrogram diagrams (Figs. 8(b)–8(i)), respectively. These images demonstrate that the carbides are composed of C and Mo. Furthermore, the car-bides are identified as Mo2C with a close-packed hexagonal structure through the calibration of the SAED pattern, as shown in Fig. 9(b). It is a hard and brittle phase which reduces the toughness of this steel.51)

During the tempering process, C atoms have enough capacity to diffuse out of the martensite, forming carbides. Among the alloying elements, the binding ability of Mo with C is relatively strong, with a significant Mo content in 300M steel. Consequently, the compounds with C and Mo are readily formed during tempering. Because the C atoms have enough time to diffuse into the matrix if two temper-ing treatment processes are performed consecutively, more Mo2C carbides are formed. The spectrogram demonstrates that the Mo2C carbides after two tempering processes tend to aggregate and expand, forming larger carbides. When DCT is introduced after the first tempering, the number of carbides is significantly decreased. The primary reason is that DCT can stabilize the existence of the C atoms in RA, inhibiting the precipitation of C atoms to form carbides with alloying elements in the subsequent heating process. As a result, the content of carbides in QTCT samples is between the QTT and QTC samples.

To better display the evolution of microstructures, a schematic describing the effects of different sequences of DCT and tempering on microstructures was established, as shown in Fig. 10. Based on the analysis above, the microstructure evolution process is speculated as follows. Many blocky RAs, a few filmy RAs, and some carbides are formed after quenching (Fig. 10(a)). When DCT is carried out directly after quenching, TMs are formed, and a little martensite formed in some blocky RAs (Fig. 10(b)). During consecutive tempering, many carbides and some martensite are formed (Figs. 10(c)–10(d)). During the QTC process, the blocky RAs are divided into several parts by newly transformed martensite, and the unconverted RAs are filmy formations (Fig. 10(e)), with a smaller amount of carbides when compared with Fig. 10(d).

According to these microstructure states a schematic of crack formation and propagation considering TMs, blocky RAs, filmy RAs, and carbides is illustrated in Fig. 11. Firstly, significant stress concentration occurs at the hard

and brittle carbides and TMs when 300M steel is subjected to external loading (Fig. 11(a)). Then, micro-cracks generate when the stress exceeds the critical value. With the external load applied, the formed micro-cracks will propagate in the matrix (Fig. 11(b)), subsequently encountering blocky RAs and filmy RAs, respectively (Figs. 11(c)–11(d)). When the micro-cracks encounter blocky RAs, strain-induced mar-tensitic transformation occurs due to the slight plastic defor-mation of the blocky RAs, further aggravating the stress concentration. As a result, secondary cracks are formed, accelerating the propagation of cracks, as shown in Figs. 11(c1)–11(c4). When the micro-cracks encounter filmy RAs, the stress concentration at the crack tip is relieved, halting the propagation of a portion of the micro-cracks, which partially hinders the formation of macro-cracks, as shown in Figs. 11(d1)–11(d2).

Combining the results of Fig. 11 and the microstructures of different treatment processes, it can be concluded that compared with other treatment processes, the microstructure produced by QTC is beneficial to the improvement in impact toughness, fracture elongation, and fatigue life. Firstly, the formation probability of micro-cracks would be reduced

Fig. 9. TEM results show (a) carbide in the QTT sample and (b) the SAED pattern of carbide in (a). (Online version in color.)

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because of the decrease in carbides and the absence of TM, limiting the possible sites for crack initiation. However, the formation of cracks cannot be prevented completely, and they will propagate in the matrix with the application of an external load after formation. The crack propaga-tion is partially hindered because unstable blocky RAs are basically eliminated and more filmy RAs are formed after treated by QTC. As a result, the fracture is delayed, leading to an improvement in impact toughness, fracture elonga-tion. Besides, the fatigue life is also significantly improved despite the slightly increase of strength owing to the favor-able microstructure.

4. Conclusions

In this study, the effects of a new DCT process (QTC) and the conventional DCT process (QCTT) on the micro-structures and mechanical properties of a typical ultrahigh strength steel were systematically investigated. Compared with the heat treatment process (QTT) and the conventional DCT process (QCTT), the impact toughness and fatigue life are simultaneously improved by the QTC process. When DCT is introduced into the heat treatment process of this medium-carbon and low-alloy ultrahigh strength steel in a conventional way (that is DCT is performed directly after quenching), the strength is not enhanced like other steels, and on the contrary, the toughness is decreased. This result is attributed to the formation of more bainite and twinned martensite due to the segregation and accumulation of car-bon atoms during QCTT process. However, when DCT is performed after tempering (namely QTC process), firstly,

the precipitation of carbide and formation of twinned mar-tensite are inhibited, decreasing the probability of crack formation. In addition, more filmy RAs with good mechani-cal stability and plasticity are transformed from blocky RAs, hindering the propagation of cracks. Combining these two aspects, the impact toughness, fracture elongation and fatigue life are simultaneously enhanced after treated by the new deep cryogenic treatment process.

AcknowledgmentsThis research was funded by the National Natural Science

Fund for Distinguished Young Scholars of China (Grant No. 51725504), the National Natural Science Foundation of China (Grant No. 51675200), and the Fundamental Research Funds for the Central Universities (Grant No. 2019kfyXMBZ030).

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Fig. 10. Schematic to illustrate the effects of different treatment processes on microstructure evolution. M: martensite; RA: retained austenite; TM: twinned martensite, M’: martensite formed after DCT, M”: martensite formed after tempering, BRA: block RA, FRA: filmy RA. (Online version in color.)

Fig. 11. Schematic to illustrate the process of crack formation and propagation. M: martensite; RA: retained austenite; TM: twinned martensite, BRA: block RA, FRA: filmy RA. (Online version in color.)

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