Effect of tempering temperature on the microstructure and properties of ultrahigh-strength stainless steel

Journal of Materials Science & Technology 35 (2019) 1240–1249

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Research article

Effect of tempering temperature on the microstructure and properties of ultrahigh-strength stainless steel Yangpeng Zhang a,b , Dongping Zhan a,∗ , Xiwei Qi b , Zhouhua Jiang a a b

School of Metallurgy, Northeastern University, Shenyang, 110819, China School of Materials Science and Engineering, Northeastern University, Shenyang, 110819, China

a r t i c l e

i n f o

Article history: Received 28 October 2018 Received in revised form 12 December 2018 Accepted 28 December 2018 Available online 22 January 2019 Keywords: Age hardening Austenite Precipitation Tempering Strengthening mechanism M7 C3

a b s t r a c t The microstructure, precipitation and mechanical properties of Ferrium S53 steel, a secondary hardening ultrahigh-strength stainless steel with 10% Cr developed by QuesTek Innovations LLC, upon tempering were studied by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD), and tensile and impact tests. Based on these results, the influence of the tempering temperature on the microstructure and properties was discussed. The results show that decomposition occurred when the retained austenite was tempered above 440 ◦ C and that the hardening peak at 482 ◦ C was caused by the joint strengthening of the precipitates and martensite transformation. Due to the high Cr content, the trigonal M7 C3 carbide precipitated when the steel was tempered at 400 ◦ C, and M7 C3 and M2 C (5–10 nm in size) coexisted when it was tempered at 482 ◦ C. When the steel was tempered at 630 ◦ C, M2 C and M23 C6 carbides precipitated, and the sizes were greater than 50 nm and 500 nm, respectively, but no M7 C3 carbide formed. When the tempering temperature was above 540 ◦ C, austenitization and large-size precipitates were the main factors affecting the strength and toughness. © 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.

1. Introduction Secondary hardening ultrahigh-strength steel (UHSS) is widely used in aeronautics and astronautics, especially in landing gear, as it has both high strength and high toughness [1]. Representative secondary hardening UHSSs include HY180, AF1410, and AerMet 100 [2–6]. HY180 was developed by the United States Steel Corporation (U.S. Steel) in 1965 [7]; its composition is Fe–0.11C–10Ni–8Co–2Cr–1Mo. In 1978, a higher-strength steel, AF1410, was developed by General Dynamics Corporation [8]; its composition is Fe–0.15C–10Ni–14Co–2Cr–1Mo. AF1410 has higher C and Co contents than HY180. AerMet 100 steel was developed in 1991 by Carpenter Technology Corporation; its composition is Fe–0.23C–11.1Ni–13.4Co–3.1Cr–1.2Mo [5,9]. All three steels have good strength and toughness, especially AerMet 100 steel. The tensile strength and fracture toughness of AerMet 100 steel are 1978 MPa and 115 MPa/mm2 , respectively [10,11]. However, compared with stainless steel, these steels have lower Cr contents and exhibit poor corrosion resistance. Failure due to stress corrosion (SCC) is very common, so a poisonous Cd protective coating is

∗ Corresponding author. E-mail address: [email protected] (D. Zhan).

needed on the surface of these steels [12]. Therefore, improving corrosion resistance has become an important research direction in the development of secondary hardening UHSS [13,14]. In 2003, QuesTek Innovations LLC developed a stainless UHSS called Ferrium S53 using a design methodology [15]. This steel has similar strength to AerMet 100 but better resistance to general corrosion and SCC [16,17]. To improve its corrosion resistance, Ferrium S53 steel contains 10 wt% Cr, which is higher than that of other UHSSs [2–4,6]. Cr is both a ferrite-stabilizing element and a strong carbide-forming element. When the Cr content is increased, the microstructure and precipitation behaviour of UHSS under different tempering and ageing processes are affected. Therefore, studying the strengthening and toughening mechanism of a secondary hardening UHSS with such a high Cr content is very important so that a steel with the best corrosion resistance, strength, and toughness can be developed. At present, only a few public reports concerning the effect of the Cr content on secondary hardening UHSS exist, and several published articles are on AF1410 steel, with Cr content up to 3% [18–20]. Seo et al. [16] studied the precipitation in Ferrium S53 steel when tempered at 500 ◦ C. Three types of carbides were found: an MC carbide containing V and Ti, an M2 C carbide rich in Mo and W, and a larger M23 C6 carbide precipitated at the grain boundaries. Except for a report by Seo et al. [16], few studies on the microstructure

https://doi.org/10.1016/j.jmst.2019.01.009 1005-0302/© 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.

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Fig. 1. (a) XRD patterns and (b) austenite contents of different samples.

and properties of Ferrium S53 steel, especially on the effects of different tempering temperatures, exist. Therefore, the relationship between the microstructure and mechanical properties is still not sufficiently understood. In the present work, the variations in the austenite content and precipitates in Ferrium S53 under different tempering processes were analysed by scanning electron microscopy (SEM), transmission electron microscopy (TEM), and X-ray powder diffraction (XRD), and the strength and impact toughness were tested. The relationship among the tempering temperature, microstructure, and mechanical properties of Ferrium S53 steel was discussed.

2. Experimental methods Ferrium S53 was produced by vacuum induction melting, followed by vacuum arc remelting. The composition (in wt%) was Fe–0.20C–14Co–5.7Ni–10Cr–2Mo–1W–0.3 V. The samples used in the heat treatment experiment were 13 mm × 20 mm × 65 mm in size. All the forged samples were homogenized at 1080 ◦ C for 1 h, cooled in air, immediately transferred into a mixture of dry ice and alcohol (−73 ◦ C) for 1 h, warmed to room temperature in air, and finally stress-relief annealed at 680 ◦ C for 8 h. After annealing, the samples were solid-solution treated at 1085 ◦ C for 1 h and then quenched by water; these samples were named SQ . These quenched samples were immediately transferred to a mixture of dry ice and alcohol (−73 ◦ C) for one hour and heated in air to room temperature. These cryogenically treated samples were named SC . The cryogenic samples were then tempered at 200 ◦ C, 300 ◦ C, 400 ◦ C, 440 ◦ C, 460 ◦ C, 482 ◦ C, 501 ◦ C, 540 ◦ C, 580 ◦ C, and 630 ◦ C and designated S200 , S300 , S400 , S440 , S460 , S482 , S501 , S540 , S580 , and S630, respectively. All samples were etched with a solution of 1.5 g CuCl2 in 33 ml HCl and 33 ml H2 O at room temperature. The etching time was 35 s, and a soaking method was used. The microstructures were first studied using ULTRA-X55 field emission SEM. Then, samples for TEM were subjected to twin-jet electropolishing with a TenuPol-5 model at −20 ◦ C and−30 ◦ C using a solid solution of 12.5% perchloric acid and 87.5% alcohol. The transmission samples were analysed by Tecnai G2 20 or JEM-2100 F TEM at 200 kV. The tensile properties of the samples were measured using a Shimadzu AG-Xplus electronic universal testing machine. The crosshead speed was 1 mm/min, and the parallel segment of the tensile samples was 30 mm × 3 mm × 2.5 mm in size. Impact tests were performed with a JBW-500 machine on V-notch Charpy impact samples (notch depth = 2 mm) with dimensions of 5 mm × 10 mm × 55 mm.

XRD measurements were carried out to determine the austenite volume fraction. Data collection was conducted using CuK␣ radiation and an angular interval of 35◦ –105◦ . A direct comparison method was used to determine the austenite volume fraction as follows [21]: V =

1 1+

R I˛ R˛ I

(1)

where V is the austenite volume fraction, and I and I˛ are the integrated intensities of the austenite and martensite peaks, respectively; R˛ and R depend on the Miller index of each phase and can be calculated [21]. The (220)˛ , (211)˛ , (200) , (220) , and (311) diffraction lines were selected to measure the integrated intensities. 3. Experimental results 3.1. XRD results Fig. 1 shows the XRD results and the retained austenite contents calculated by the direct comparison method. According to Fig. 1(b), the austenite content of quenched sample SQ is the highest, reaching 35.5%, while the austenite content of SC decreases to 22.6% after cryogenic treatment. When the steel is tempered at 200 ◦ C, the austenite content is the same as that of SC , 22.6%. The austenite content increases to 27.8% at 300 ◦ C tempering, decreases to 21.8% at 400 ◦ C, and increases to 28.9% at 440 ◦ C. The austenite content then decreases sharply when the tempering temperature exceeds 440 ◦ C and reaches a minimum between 501 ◦ C and 540 ◦ C, approximately 3%. The austenite content then increases rapidly when the tempering temperature is increased further, reaching 30.8% at a tempering temperature of 630 ◦ C. 3.2. SEM results The microstructures of S200 , S400 , S460 , S482 , S540 , and S630 after tempering are lath martensite, as shown in Fig. 2. After 630 ◦ C tempering (high temperature), the samples are quite different from the other samples, as shown in Fig. 2(f): etched fine phases appear in the sample after corrosion, which are likely to be austenite based on the observed variation trend of the retained austenite. 3.3. TEM results Fig. 3 shows TEM images of quenched sample SQ . The irregular black stripes are austenite and 200 nm in width. As shown in Fig. 3(a), the martensite is sandwiched between austenite and is

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Fig. 2. Microstructures of different samples: (a) S200 ; (b) S400 ; (c) S460 ; (d) S482 ; (e) S540 ; (f) S630 .

Fig. 3. (a) Bright field image of austenite, (b) dark field image of austenite and (c) lath martensite morphology of sample SQ .

blocky in shape; in addition to this martensite, lath martensite is also found, as shown in Fig. 3(c). The upper right corner of Fig. 3(c) shows the calibration result of the lath martensite diffraction spot. Fig. 4 shows a TEM image of sample SC . The microstructure of the steel is mainly martensite. In addition, some short-range ordered regions exist in the samples, as shown in Fig. 4(c). Fig. 5 shows the microstructure and precipitates of sample S400 , which was tempered at 400 ◦ C for 6 h. In this sample, the main microstructure is martensite, and the main secondary phase is needle-like M7 C3 carbide, as shown in Fig. 5(a). A high-resolution transmission electron microscopy (HRTEM) image of the M7 C3 carbide is shown in Fig. 5(b), and the width is approximately 10 nm.

The diffraction spot of the carbide, obtained by fast Fourier transform (FFT) of the red area in Fig. 5(b), is shown in the upper right of Fig. 5(b). The calibration results prove that this area is M7 C3 carbide with an orthorhombic structure, and the lattice constants are a = 0.454 nm, b = 0.6879 nm, c = 1.1942 nm, and ˛ = ˇ =  = 90◦ . In addition to needle-like carbides, another smaller carbide, which has no obvious interface with the matrix, exists, as shown in FFT3 of Fig. 5(c). The diffraction spots of matrix region FFT2 and carbide region FFT3 obtained by FFT are shown in Fig. 5(d) and (e), respectively. The diffraction calibration results show that these small carbides are also M7 C3 and have the same orthorhombic structure as the needle-like M7 C3 . The martensite and M7 C3 carbide

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Fig. 4. Microstructure of sample SC : (a) martensite matrix and diffraction calibration; (b) HRTEM image of martensite and diffraction calibration; (c) short-range order regions.

Fig. 5. Microstructure and precipitates in sample S400 : (a) martensite and needle-like M7 C3 ; (b) HRTEM image of M7 C3 and diffraction calibration; (c) HRTEM images of small M7 C3 carbides; (d, e) FFT and diffraction calibration of the martensite matrix and small M7 C3 .

have the following orientation relationship: [111]˛ //[100]M 7C3 and (110)˛ //(031)M 7C3 . Fig. 6 shows the microstructure of samples after tempering at 482 ◦ C. According to Fig. 6(a), lath martensite is the major microstructure of sample S482 , and no obvious large-size secondary phase is found in the field of view. Inside some laths, some nanocrystalline martensite twins exist. The twins seen in the red frame of Fig. 6(b) have been amplified to analyse the nanosize precipitates and microstructures, as shown in Fig. 7. As shown in Fig. 7(a), some dislocations and fine precipitates exist in the twins. Fig. 7(b) presents the electron diffraction spot and calibration result at this position. The grey mesh was drawn from symmetrical electron diffraction spots, which are the electron diffraction patterns (EDPs) of the martensite twins. In addition to the EDP of the matrix, two types of carbide EDPs are shown in Fig. 7(b), corresponding to the M7 C3 carbide, calibrated by the yellow parallelogram, and the M2 C carbide, calibrated by the white parallelogram. The lattice structure of M7 C3 is the same as that of the M7 C3 in sample S400 , both of which are orthorhombic, and the lattice constants are also the same as the aforementioned ones. The lattice structure of the M2 C carbide is hexagonal, with lattice constants of a = b = 0.2994 nm, c = 0.4772 nm, and ˛ = ˇ = 90,  = 120, which are the same as those previously reported [2,22,23]. HRTEM

images of some black precipitates in Fig. 7(a) are shown in Fig. 7(c). The diffraction spots of the FFT4 region in Fig. 7(c) obtained by FFT are shown in Fig. 7(d). The calibration result indicates that the black phase is M7 C3 , and the Z axis is [010]. Fig. 8 shows TEM images of sample S630 , which was tempered at 630 ◦ C for 6 h. Fig. 8(a) shows that plenty of lath martensite still exists in the steel, but part of the lath boundary has been dissolved. In addition, secondary phases appear in the lath martensite. Unlike the samples tempered at lower temperatures, a large amount of massive austenite phase appears in the steel, as seen by the black phases in Fig. 8(b). The length of these austenite phases can reach several microns, and the width is greater than 200 nm. Fig. 8(c) shows the magnification and diffraction calibration results of these black austenite blocks, which coincide with the diffraction spots of austenite, and Fig. 8(d) shows the dark field image of the corresponding position. Fig. 8(c) and (d) shows that the austenite grains are mostly approximately 500 nm in size, but the secondary phases inside the martensite matrix are smaller than 100 nm. Fig. 9 shows the mapping image of sample S630 . According to the bright field image, many phases of different sizes and shapes exist in the sample. The mapping results show that the largest phases are Ni enriched compared to the matrix, which corresponds to the austenite described above. In addition, the Cr content is higher in

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Fig. 6. Microstructure of sample S482 : (a) lath martensite; (b) martensite twins.

Fig. 7. Carbides in sample S482 : (a) carbides in martensite twins; (b) EDP and calibration of M7 C3 and M2 C carbides; (c) HRTEM image of M7 C3 ; (d) calibration of M7 C3 .

Fig. 8. Microstructure of sample S630 : (a) lath martensite; (b) large-size austenite; (c) bright field image of austenite; (d) dark field image of austenite.

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Fig. 9. Mapping results of sample S630 (BF: bright field image). Table 1 Composition of secondary phases in sample S630 . Position

Fe

Co

Cr

Ni

Mo

W

V

Experimental steel  1  2  3  4  5  6

67 73.65 56.90 41.71 55.2 39.23 49.68

14 15.92 13.43 9.82 12.08 5.76 11.14

10 6.38 14.1 32.90 12.34 39.10 11.35

5.7 2.62 10.05 2.30 2.83 0.96 2.93

2 0.53 3.18 8.37 12.15 7.88 18.10

1 0.64 1.59 18 5.89 6.10 6.54

0.3 0.25 0.45 0.73 0.20 0.98 0.25

in Table 1) were selected for diffraction calibration and dark field image analysis. The results are shown in Fig. 10. Fig. 10(a) shows that the size of precipitate  5 is greater than 500 nm, and the diffraction calibration result in Fig. 10b indicates that it is carbide M23 C6 . Precipitate , 6 shown in Fig. 10(c), is approximately 50–100 nm in size. The diffraction calibration result in Fig. 10(d) shows that the phase is an M2 C carbide with a hexagonal structure, and the lattice constants are the same as those of the M2 C carbides in sample S482 , a = b = 0.2994 nm, c = 0.4772 nm, and ˛ = ˇ = 90◦ , ␥ = 120◦ . 3.4. Mechanical properties

the austenite than in the matrix. In the smaller secondary phases, the more distinctly enriched elements are Cr and Mo, and the most depleted elements are Fe and Co. This phenomenon occurs because Cr and Mo are strong carbide-forming elements, carbides rich in Cr and Mo are easily formed, and other elements are discharged from the carbide positions. We also used EDS to quantitatively analyse multiple positions in the field of view of Fig. 9. The results show that the phases are divided into four categories according to their components, and the typical phases are marked as , 1 , 2 , 3 and  4 in the bright field image of Fig. 9. The results of the EDS quantitative analysis are shown in Table 1. According to Table 1, the contents of the alloying elements Cr, Ni, Mo, and W in the martensite matrix  1 are lower than those in the steel composition due to the precipitation behaviour of other secondary phases, including austenite phases  2 and carbides  3 and . 2 is 10.05%, which 4 The Ni content in the austenite phase  is greater than 5.7%, the Ni content of the experimental steel. It is also higher than that of all the other phases, which proves that Ni is enriched in the austenite phase. The phase that has the highest Cr content (32.9%) is phase , 3 with a size of 150 nm, and the Mo, W, and V contents in this phase are also significantly higher than those in the steel. In addition, although the Fe and Co contents in phase  3 are higher than those in the other phases, they are lower than those in the test steel. Another typical precipitate is phase , 4 which is smaller than phase  3 and has higher Mo and W contents—12.15% and 5.89%, respectively. To study the structure of phases  3 and , 4 two other precipitates with similar compositions (phases  6 5 and 

The tensile test results of all samples at room temperature are shown in Fig. 11. As shown in Fig. 11(a)-(c), the tensile strength (TS) and yield strength (YS) of the quenched sample are very low, especially the YS, which is only 546 MPa. After cryogenic treatment, the YS and TS clearly increase. YS increases from 546 MPa to 1177 MPa, and TS reaches 1815 MPa. The TS and YS of the samples tempered at 200 ◦ C and 300 ◦ C decrease little compared with sample SC , and the TS and YS of sample S300 are 1460 MPa and 1095 MPa, respectively. When the tempering temperature is above 300 ◦ C, the TS and YS both begin to increase rapidly and reach their highest values in sample S482 , 1945 MPa and 1543 MPa, respectively. When the tempering temperature is raised above 482 ◦ C, the YS and TS decrease as the temperature is continuously increased. The TS and YS decrease to 1233 MPa and 920 MPa, respectively, when the steel is tempered at 630 ◦ C. The change of elongation with tempering temperature is fundamentally consistent with that of the impact toughness but is contrary to that of the strength, as shown in Fig. 11(a). Both the elongation and impact toughness of SQ are the highest among all the samples; the elongation is 23.3%, and the impact energy is 64 J. After cryogenic treatment, the elongation and impact toughness decrease significantly, especially the impact energy, which is reduced to 22.5 J, a drop of 65%. However, when the SC sample is tempered at 200 ◦ C, the elongation and impact toughness increase to the levels of the quenched sample (SQ ), 21.6% and 63 J, respectively. When the tempering temperature is above 200 ◦ C and increased further, both properties maintain a downward trend, and

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Fig. 10. Carbides in sample S630 : (a) M23 C6 -type carbide; (b) calibration of M23 C6 ; (c) M2 C-type carbide; (d) calibration of M2 C.

Fig. 11. Mechanical properties of samples: (a) strain-stress curves; (b) TS and YS; (c) elongation and impact toughness.

the rate of decrease increases significantly when the tempering temperature exceeds 400 ◦ C. The elongation and impact toughness reach minimum values at 540 ◦ C and 500 ◦ C, respectively, that is, 9.5% and 17.5 J, respectively. However, when the tempering temperature is further increased, the elongation (EL) and impact toughness (Akv ) begin to increase rapidly once more and reach 13.4% and 3.5 J at 630 ◦ C, respectively. 4. Discussion 4.1. Effect of tempering temperature on austenite The austenite content is an important factor that affects the properties of martensitic steel [4,24]. A small amount of austenite, especially reversed austenite, can improve the toughness of steel [5,9,25], but excessive austenite can cause a significant decrease in strength [11]. Fig. 12 shows the dilatometric curve of Ferrium S53 steel measured by the phase transformation tester Formastor-FII. The austenite transformation start temperature Ac1 , the austenite transformation end temperature Ac3, and the martensite transformation start temperature Ms of the steel are 590 ◦ C, 780 ◦ C, and 100 ◦ C, respectively. The Ms temperature of the steel, measured by a thermal expansion experiment, is consistent with the value given by QuesTek Innovations LLC [17], which is only 100 ◦ C, very low for the marten-

Fig. 12. Dilatometric curve of Ferrium S53 steel.

sitic transformation start temperature. The Ms temperature is an important factor affecting the amount of retained austenite, and Cr can reduce the Ms value, resulting in 35.5% austenite after quenching [26,27]. To improve the transformation ratio of martensite, a subzero treatment or a cryogenic treatment is usually used to increase the degree of undercooling, which enhances the driving

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forces of martensite transformation and reduces the amount of retained austenite [9,22,28]· Therefore, as expected, the calculation result in Fig. 2 shows that the austenit ze content decreases to 22.6% after cryogenic treatment. The change in austenite during tempering is mainly affected by the decomposition of the retained austenite and the formation of reversed austenite [29–32]. When the cryogenically treated sample is tempered at 200–440 ◦ C, the temperature is much lower than the Ac1 temperature [33], so reversed austenite cannot form. Therefore, if the retained austenite does not undergo decomposition and transformation, the austenite content does not change very much compared to the cryogenic sample. Internal stress and the Cottrell atmosphere usually resist the transformation of retained austenite [34,35]. When the tempering temperature is sufficiently high, the internal stress and C and N in the Cottrell atmosphere will be released; therefore, during cooling, martensitic transformation occurs, which is also called conditioning of retained austenite [36,37]. As shown in Fig. 1(b), this critical temperature is approximately 440 ◦ C. When the tempering temperature is between 440 ◦ C and 540 ◦ C, the austenite content is significantly lower than that of the cryogenic sample. The reason is that, on the one hand, the conditioning of retained austenite becomes faster with increasing temperature [34–37]. On the other hand, austenite instability is caused by the carbon concentration reduction due to carbide precipitation. When the temperature is low, the diffusion of C and other elements is difficult, and carbides do not easily precipitate. When the temperature is sufficiently high, the diffusion rate of C increases, and carbides precipitate, resulting in a decrease of carbon in the retained austenite matrix and austenite instability. The detailed precipitation behaviour of carbides is described in part 4.2. The austenite content in sample S580 is lower than that in sample SC but higher than that in sample S540 . First, the decomposition of retained austenite still occurs, and the retained austenite content is still lower than that of the SC sample. However, because 580 ◦ C is close to the critical temperature Ac1 , nickel-rich martensite may have already transformed to austenite, and a large amount of reversed austenite is formed, resulting in a higher austenite content in S580 than in S540 . When the temperature is increased to 630 ◦ C, which is 40 ◦ C higher than the Ac1 point, austenitization will occur in large quantities. Although some retained austenite is still transformed to martensite again during cooling, the total content is increased significantly compared to sample SC and is close to that of sample SQ . As shown in Fig. 8, ellipsoidal or strip-shaped austenite grains of different sizes exist. The mapping results in Fig. 9 and Table 1 show that this part of the austenite is obviously enriched in both Ni and Cr. 4.2. Effect of tempering temperature on precipitation Precipitates are another important factor that affects the strength and toughness of UHSS [1,38,39]. When the solid solution is treated at 1085 ◦ C, Ferrium S53 steel is completely austenitized. After one hour of solid-solution treatment, various types of carbides produced in the forging and annealing stage are dissolved, and precipitates will precipitate during tempering [6]. Cryogenic treatment is mainly used to increase the undercooling of retained austenite and promote martensite transformation. Some research has reported that small amounts of cementite can precipitate and act as a nucleation core for secondary phases during tempering [40–42]. In this study, no cementite with obvious interfaces is observed in sample SC , but as shown in Fig. 4(c), some short-range ordered structures exist in the matrix, with a size of only 2 nm. These short-range ordered structures may also become nucleation cores, but no direct evidence for this has been found at present. Many papers [1,11,22,23,39,43] have reported that when CoNi secondary hardening UHSS is tempered at temperatures near

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Fig. 13. EDS result of the M7 C3 carbide in sample S482 .

482 ◦ C, the main precipitate phase is M2 C-type carbide and is rich in Mo and W. In this experiment, M2 C-type carbide was also found in sample S482 , but the calibration spot was very weak, which means that this type of carbide is small in size or there is little of it. Unlike in Co-Ni steel, in the experimental Ferrium S53 steel, orthorhombic M7 C3 carbides with lattice constants of a = 0.445 nm, b = 0.6879 nm, c = 1.1942 nm, and ˛ = ˇ =  = 90◦ are found in samples S400 and S482 . Carbides directly converted from cementite in situ are usually larger in size and do not cohere with the matrix, while carbides precipitated from the martensite matrix are usually smaller in size and are coherent or semicoherent with the matrix [44]. The M7 C3 carbides observed in sample S400 are mostly thin section shaped. The thin section width is only approximately 10 nm, martensite exists between the flakes, and the M7 C3 carbide observed at 482 ◦ C is also only approximately 10 nm in size. Such a small-size M7 C3 carbide should be directly formed from the matrix, and in the S482 sample, the carbide is indeed in the twin, which matches the characteristics of direct nucleation. Moreover, as the calibration results in Fig. 5 show, after tempering at 400 ◦ C, the M7 C3 precipitates have the following orientation relationship with the matrix: (101)˛ //(230)M 7C3 and [110]˛ //[001]M 7C3 . Therefore, the M7 C3 carbide is directly precipitated from the martensite matrix and maintains a coherent relationship with the matrix. If prior precipitated cementite exists, these cementite phases should have dissolved into the matrix before the precipitation of M7 C3 . The M7 C3 carbides found in the samples tempered at 400 ◦ C and 482 ◦ C are small in size and can play a very influential strengthening role without weakening the toughness and plasticity. For sample S482 , one EDS result of a precipitate is shown in Fig. 13. Since the size of the precipitated phase is small, the EDS result includes the signal of the matrix. However, compared with the peak corresponding to the composition of the experimental steel, the Cr peak is obviously higher. This result means that these precipitates are rich in Cr, and many other scientists have also found Cr-rich M7 C3 carbides in 9–12 wt% Cr steel [45–48]. Wang [47] observed carbides in Fe-10Cr0.15C (wt%) martensitic steel and found similar orthogonal M7 C3 carbides. The precipitation temperature of M7 C3 is 500–700 ◦ C, and when the temperature is above 700 ◦ C, carbide M23 C6 will precipitate. Therefore, unlike Co-Ni steel, due to the addition of 10% Cr, Ferrium S53 steel is more likely to precipitate M7 C3 -type carbides as a secondary hardening phase during tempering at 400 ◦ C and 482 ◦ C. When samples are tempered at 630 ◦ C, the coarsening rate of carbides is significantly increased due to the high temperature. M23 C6 carbides rich in Cr (more than 30%) appear, and their size

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can exceed 500 nm, which is disadvantageous to the properties of the steel. However, no M7 C3 carbide is found in this sample, similar to the steel annealed at 680 ◦ C [6]. In addition to Cr-rich M23 C6 carbides, Mo-rich and W-rich M2 C carbides have been found in the S630 sample. The size of the M2 C carbide is approximately 50 nm, which is much larger than the size of the M2 C carbide observed in sample S482 , but it still has a good strengthening effect on the steel. 4.3. Relationship between temperature, microstructure, and properties For sample SQ , the YS is only 546 MPa, which is much lower than that of other samples, but the TS is as high as 1648 MPa. The XRD and TEM analysis results show that there is 35.5% retained austenite in sample SQ , and the strength of austenite is much lower than that of martensite. In addition, there is a lack of precipitates to prevent dislocation movement, so plastic deformation begins when the stress exceeds 546 MPa. Under continuous deformation, retained austenite transforms into martensite due to deformation-induced martensitic transformation, which improves the resistance to continuous deformation and the strength. Due to the delay of necking, the ductility is also improved. Therefore, compared with the other samples, the uniform deformation stage of SQ is the longest. Compared with sample SQ , the content of retained austenite after cryogenic treatment decreased to 22.6%. Martensitic transformation means more residual stresses (especially microresidual stress) [49–51] and the formation of cementite or an ordered phase, as shown in Fig. 3. Therefore, dislocation glide is more difficult, and the YS is improved. The TSs after tempering at 200 ◦ C and 300 ◦ C are 1565 MPa and 1460 MPa, respectively, and the YSs are 1153 MPa and 1095 MPa, respectively. Compared to the values of 1815 MPa and 1197 MPa for sample SC , both the TS and the YS are reduced. The retained austenite content of S200 is the same as that of SC , 22.6%. Therefore, the austenite content is not the main reason for the difference in the mechanical properties. After quenching and cryogenic treatment, sample SC has a very high degree of lattice distortion and internal stress [40]. However, after tempering, the lattice distortion in the sample partially recovers, and the internal stress decreases accordingly [51–53]; thus, the strengthening effect caused by microresidual stress will also disappear. In addition, the dislocation density and lattice distortion increased by cryogenic treatment also decrease, which leads to a decrease in strength. Therefore, when steel is tempered below 300 ◦ C, the release of microresidual stress and the change in dislocation density are factors affecting its strength and toughness. In addition to the above two factors, the partitioning of carbon is also an important factor affecting the strength and toughness. Although the steel is tempered at low temperature, 200 ◦ C is higher than Ms . Carbon will diffuse from martensite to austenite, which improves the stability of retained austenite and reduces the carbon content of martensite, resulting in a significant increase in ductility. When the cryogenically treated steel is tempered at 400 ◦ C, the TS decreases from 1815 MPa to 1640 MPa, while the YS increases from 1197 MPa to 1281 MPa. Compared with sample S300 , the TS and YS increase 180 MPa and 84 MPa, respectively. According to the classical Orowan formula for second-phase strengthening, carbides mainly increase the YS [54–56]. As mentioned in part 4.2, M7 C3 -type carbides are found in S400 ; therefore, the M7 C3 carbide improves the YS of S400 . When the steel is tempered in the range of 400–482 ◦ C, the TS continues to increase and reaches a peak of 1945 MPa at 482 ◦ C, and the YS reaches a peak of 1543 MPa. As discussed in Section 4.2, the retained austenite content of the samples decreases rapidly over this temperature range, and only 9.3% austenite is contained in sample S482 . An increasing amount of martensite provides continuous enhancement of both the YS and

TS of the samples. Precipitation strengthening is another important strengthening factor. M7 C3 carbides and M2 C carbides are precipitated in sample S482 . The secondary hardening peak at 482 ◦ C is caused by the joint strengthening of the precipitates and martensite transformation. The variation tendencies of the impact toughness and elongation are the same as those of the retained austenite, which means that the decomposition of the retained austenite has the most obvious effect on the toughness and plasticity of the steel. A critical tempering temperature of 482 ◦ C is found. When the temperature is above 482 ◦ C, the TS and YS decrease continuously with increasing temperature. At the highest tempering temperature of this experiment, 630 ◦ C, the YS decreases to 920 MPa, and the TS decreases to 1233 MPa. When the steel is tempered at 482 ◦ C, 501 ◦ C, and 540 ◦ C, its austenite contents are 9.3%, 3.2%, and 3.2%, respectively. From the previous analysis, the decrease in the austenite content will lead to increased strength, but in these three samples, as the TS and YS continue to decline, the reduction in YS is more obvious. According to many reports, 482 ◦ C is the best tempering temperature for secondary hardening steel, and if the temperature is further increased, then the precipitates and martensite begin to coarsen rapidly. Martensite coarsening will reduce the grain boundary strengthening and the toughness of the steel. Large-size precipitates will lose coherence with the matrix, and the coherent strain strengthening will disappear. In addition, due to the coarsening of precipitates, the spacing between precipitates increases. According to the Orowan formula [54–56], the secondphase strengthening will become increasingly weaker. In addition to decreasing the strength of the steel, large-size precipitates can lead to microcracks and reduce the toughness and plasticity. The growth of precipitates requires carbon and alloying elements, such as Cr, Mo, W and V, but these are also the elements that play a solid-solution strengthening role. When the content of carbon and alloying elements in martensite decreases, the strengthening caused by the solid solution decreases. Above all, the coarsening of precipitates and martensite is the main reason for the change in toughness when the steel is tempered between 482 ◦ C and 540 ◦ C. When the sample is tempered at temperatures above 540 ◦ C, the temperature is close to or exceeds the Ac1 temperature, austenitization begins, and austenite grains are formed, as shown in Fig. 8. These austenite grains enhance the continuous deformation capacity of the steel and improve its toughness and plasticity while simultaneously reducing the strength. In sample S630 , many precipitates whose size is greater than 500 nm exist. When the sample is under loading, a large number of dislocations will pile up around them, and microcracks will occur preferentially in these areas, which will lead to material failure. From the above analysis, the strength and toughness of steel tempered above 540 ◦ C are mainly affected by the austenitization and large-size precipitates.

5. Conclusions The change in the microstructure and precipitates of UHSS with high Cr content upon different heat treatments was studied in this paper. The following three meaningful conclusions were obtained:

(1) Quenched samples contain approximately 35.5% austenite, which can be reduced to 22.6% after cryogenic treatment. Decomposition of retained austenite occurs when the tempering temperature is above 440 ◦ C, and the lowest austenite content of approximately 3.2% is obtained when tempering is performed between 500 ◦ C and 540 ◦ C. The martensite to austenite transformation occurs when the sample is tempered above 580 ◦ C, and the newly formed austenite is rich in Ni and Cr.

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(2) Trigonal M7 C3 carbide precipitates when Ferrium S53 steel is tempered at 400 ◦ C, and M7 C3 and M2 C precipitate at 482 ◦ C with a size of 5–10 nm. M2 C carbides with a size of 50 nm and M23 C6 carbides with a size of 500 nm precipitate at 630 ◦ C. (3) Internal stress release and dislocation density decreases affect the strength and toughness most when the tempering temperature is below 300 ◦ C. The hardening peak at 482 ◦ C is co-strengthened by precipitates and martensite transformation. Coarsening of precipitates and martensite causes a decrease in strength and toughness when tempered at 482–540 ◦ C. When the tempering temperature is above 540 ◦ C, austenitization and large-size precipitates are the main factors that affect the strength and toughness. Acknowledgments This work was supported financially by the National Natural Science Foundation of China (Nos. 51874081 and 51574063) and the Fundamental Research Funds for the Central Universities (Nos. N152306001 and N150204012). References [1] Z.B. Jiao, J.H. Luan, M.K. Miller, Y.W. Chung, C.T. Liu, Mater. Today 20 (2017) 142–154. [2] R. Veerababu, R. Balamuralikrishnan, K. Muraleedharan, M. Srinivas, Metall. Mater. Trans. A 46 (2015) 2455–2468. [3] K.S. Cho, H.S. Sim, J.H. Kim, J.H. Choi, K.B. Lee, H.R. Yang, H. Kwon, Mater. Charact. 59 (2008) 786–793. [4] Y.P. Zhang, D.P. Zhan, X.W. Qi, Z.H. Jiang, H.S. Zhang, Mater. Sci. Eng. A 698 (2017) 152–161. [5] Y.P. Zhang, D.P. Zhan, X.W. Qi, Z.H. Jiang, Mater. Charact. 144 (2018) 393–399. [6] Y.P. Zhang, D.P. Zhan, X.W. Qi, Z.H. Jiang, Mater. Sci. Eng. A 730 (2018) 41–49. [7] M. Krenzke, K. Hom, J. Proffitt, Potential Hull Structures for Reduce and Search Vehicles of the Deep-Submergence Systems Project, 2019, 9th January https://pdfs.semanticscholar.org/c6d6/ f867f9564dd6b48d052b8749aa78fae79d7d.pdf. [8] C.D. Little, P.M. Machmeier, High Strength Fracture Resistant Weldable Steels, US Patent, No. US4076525, 1978. [9] C.C. Wang, C. Zhang, Z.G. Yang, Micron 67 (2014) 112–116. [10] R. Veerababu, R. Balamuralikrishnan, K. Muraleedharan, M. Srinivas, Metall. Mater. Trans. A 39 (2008) 1486–1495. [11] X.H. Shi, W.D. Zeng, Q.Y. Zhao, W.W. Peng, C. Kang, J. Alloys. Compd. 679 (2016) 184–190. [12] G.L. Pioszak, R.P. Gangloff, Metall. Mater. Trans. A 48 (2017) 4025–4045. [13] M.N. Rao, M.K. Mohan, P.U.M. Reddy, Corros. Sci. 51 (2009) 1645–1650. [14] L.W. Tsay, M.Y. Chi, Y.F. Wu, J.K. Wu, D.Y. Lin, Corros. Sci. 48 (2006) 1926–1938. [15] G.B. Olson, Science 277 (1997) 1237–1242. [16] J.Y. Seo, S.K. Park, H. Kwon, K.S. Cho, Metall. Mater. Trans. A 48 (2017) 4477–4485.

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