Effect of high magnetic field on the microstructure evolution and mechanical properties of M50 bearing steel during tempering

Journal Pre-proof Effect of high magnetic field on the microstructure evolution and mechanical properties of M50 bearing steel during tempering Feng Wang, Dongsheng Qian, Lin Hua, Huajie Mao, Lechun Xie, Xinda Song, Zhaohua Dong PII:

S0921-5093(19)31409-1

DOI:

https://doi.org/10.1016/j.msea.2019.138623

Reference:

MSA 138623

To appear in:

Materials Science & Engineering A

Received Date: 19 September 2019 Revised Date:

29 October 2019

Accepted Date: 31 October 2019

Please cite this article as: F. Wang, D. Qian, L. Hua, H. Mao, L. Xie, X. Song, Z. Dong, Effect of high magnetic field on the microstructure evolution and mechanical properties of M50 bearing steel during tempering, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/j.msea.2019.138623. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Effect of high magnetic field on the microstructure evolution and mechanical properties of M50 bearing steel during tempering Feng Wang a, b, c, Dongsheng Qian a, b, c, ∗, Lin Huaa, b, c,∗∗, Huajie Mao a, b, c, Lechun Xie b, c, Xinda Song a, b, c, Zhaohua Dong a, b, c a. School of Materials Science and Engineering, Wuhan University of Technology, Wuhan 430070, China; b. Hubei key Laboratory of Advanced Technology for Automotive Components, Wuhan 430070, China; c. Hubei Engineering Research Center for Green Precision Material Forming, Wuhan 430070, China;

Abstract: The effect of high magnetic field (12 T) on the microstructure evolution and mechanical properties during tempering of M50 bearing steel have been investigated for evaluating the application prospect of magnetic field in steels. Microstructural observation shows that the high magnetic field can obviously accelerate the decomposition of retained austenite (RA) during the moderate tempering (at 300 or 380 °C), which could be ascribed to the increased dislocation in RA by the magnetic field-induced deformation of martensite. Meanwhile, the cementite precipitation is also increased by the applied high magnetic field because of the higher nucleation rate. Moreover, at higher tempering temperature of 530 °C, the Fe content in the precipitated carbides increases after applying high magnetic field, which is related to the different magnetization of carbides caused by various element content. Characterization of mechanical properties indicates that the tempered martensite embrittlement is significantly intensified during tempering at 380 °C, which is resulted from the accelerated RA decomposition and improved cementite precipitation under high magnetic field. When tempered at 530 °C in high magnetic field, the dislocation recovery is retarded and then the Vickers hardness and ultimate tensile strength (UTS) are thus improved by dislocation strengthening. This finding suggests the potential of high magnetic field (12T) in optimizing the properties of M50 bearing steel during tempering (530 °C). Key words: High magnetic field; Microstructure; Carbide; Retained austenite; Mechanical

∗ ∗∗

Corresponding author: Dongsheng Qian; Prof.; Tel: +86 27 87505177; E-mail: [email protected] Corresponding author: Lin Hua; Prof.; Tel: +86 27 87168391; E-mail: [email protected]

1

property; M50 bearing steel

1. Introduction The magnetic field has been widely introduced to phase transformations of steels for the purpose of microstructure modification. Because of the different magnetic properties of parent and product phases in steels, the application of external magnetic field has a considerable impact on the thermodynamic and kinetic of the phase transformation [1, 2]. Therefore, the effects of magnetic field on various phase transformation, including austenite [3], martensite [4, 5], bainite [6, 7], ferrite [8, 9] and pearlite [10-12] transformation, have been investigated extensively. In particular, the magnetic-field-induced precipitation behaviors during martensitic decomposition draw increasing attention these years [13-16]. For instance, Zhou et al. [13] applied a magnetic field during the high temperature tempering of the Fe–C–Mo alloy, indicating that (Fe, Mo)6C alloy carbides were preferentially precipitated instead of Fe3C, (Fe, Mo)2C and (Fe, Mo)3C carbides. Hou and Wu [14] further investigated the effect of high magnetic field on alloy carbide precipitation of 2.25Cr-Mo steel during tempering at 200 °C. They found that the higher Fe atom content in the M23C6 carbide caused a great increase in the magnetic moment, thereby decreasing the magnetic Gibbs free energy of this carbide. In a high Chromium-containing steel, the high magnetic field has been proved to contribute to the nucleation of (Fe, Cr)3C carbide and thus increases the number density of precipitates [15]. In addition, the decomposition behavior of retained austenite (RA) under a high magnetic field attracts much research interest [17-19]. Kaletina [17] reported that the magnetic field led to a considerable increase in the martensitic transformation temperature and hence caused additional decomposition of RA. The works of Fokina [18] indicated that the introduction of high magnetic field could facilitate the transformation of martensite at a constant temperature, which may 2

further affect the hardness and the dimensions of machine parts. However, the mechanism explaining the magnetic field-induced transformation was not provided. By summarizing the existing researches, it can be found that the studies investigating how magnetic field affects the mechanical properties are still scare, though plenty of researches have focused on the microstructure modification caused by magnetic field. M50 bearing steel is widely used in the aerospace industry as main shaft bearing in gas-turbine engines due to its excellent elevated temperature performance, which strongly relies on secondary hardening during subsequent tempering [20]. However, the volume fraction of RA in M50 bearing steel is often less than 5% after traditional tempering, leading to a poor ductility and toughness [21]. It is thus known that the carbide precipitation and RA decomposition behaviors during tempering play a major role in determining the mechanical properties of M50 bearing steel. Considering the magnetic field-induced behaviors of carbide precipitation and RA decomposition, it is expected that the mechanical properties could be influenced by high magnetic field. However, there are few systematic works concerning about the influence of high magnetic field on the microstructure evolution and mechanical properties during tempering of secondary hardening steel. In this research, the non-isothermal decomposition of RA and precipitation of carbides with different tempering parameters without and with high magnetic field (12 T) in M50 bearing steel have been systematically investigated. Furthermore, in order to clarify the relationship between high magnetic field, microstructure and mechanical properties during the tempering of M50 bearing steel, the hardness and tensile property were analyzed in detail considering the effect of high magnetic field. 2. Experimental details

3

2.1 Material and processing The chemical composition of M50 bearing steel is presented in Table 1. The initial material was received as spheroidise annealed bar and machined into cylindrical specimens with dimensions of 25 mm in height and 14 mm in diameter. The specimens were first austenitized at 1090 °C for 30 min in a WZGQ-30 vacuum furnace. A mixed microstructure consisting quenched martensite and retained austenite was then obtained by oil quenching at 60 °C. The volume fraction of RA in the as-quenched specimen is determined to be 26.5 vol. %. After quenching, using a vacuum resistance furnace installed with a superconducting magnet, some specimens were tempered at 300 and 380 °C for 10 and 60 min, and then the rest of specimens were tempered at 530 °C for 120 min. During the tempering, the as-quenched specimens were first placed in the central region (zero magnetic force) with their longitudinal direction parallel to the axis of the magnet. After the strength of magnetic field reached the set value (0 or 12 T), the resistance furnace was heated at a rate of 5 °C/min and isothermally held at the set temperature in vacuum. Finally, all specimens were furnace-cooled to room temperature. Table 1 Chemical compositions of M50 bearing steel (wt.%). C

Cr

Mo

V

Mn

Si

W

Fe

0.8~0.85

4~4.25

4~4.5

0.9~1.1

0.15~0.35

≤0.25

≤0.25

Bal.

2.2 Microstructure characterization The microstructure evolution of the specimens tempered without and with high magnetic field (12 T) was characterizated by a field emission scanning electron microscope (FESEM, FEI Quanta 450) and a transmission electron microscope (TEM, JEOL 2100F). The specimens for SEM were mechanical polished and then etched in an alcohol solution containing 4% nitric acid (volume fraction) for 30 s. The thin foil specimens for TEM experiments were prepared on a 4

twin-jet polisher at 20V using a solution of 10 % perchloric acid and 90 % acetic acid. Additionally, a Netzsch DIL 402C dilatometer was used to investigate the effect of magnetic field on the microstructure transformation during tempering. The dilatometry specimens were cut into a size of φ 4 mm × 20 mm and then heated up to a temperature of 530 °C at the heating rate of 5 °C /min. For observing the effect of high magnetic field on the carbide precipitation behaviors, the carbides in specimens were extracted by a solution consisting of 8 % HCl and 2 % tartaric acid in methanol under the condition of the temperature at 0 ~ 5 °C and current density about 50 mA/cm2. After that, the extracted carbides powder was analyzed by means of a vibrating sample magnetometer (Squid-VSM). The solutions collected after electrolytic extraction were thus free of carbides, and then, an inductively coupled plasma optical emission spectroscopy (ICP-OES) was employed to evaluate the chemical elements of the matrix. 2.3 X-ray diffraction analysis The specimens for X-ray diffraction (XRD) analysis were polished mechanically and electrolytically, and then the X-ray diffraction experiments were performed with a scanning speed of 1°/ min on a Rigaku D/MAX-RB diffraction analyser with Cu Kα1 radiation at 12 kW. The volume fraction of RA was determined by the following equations [22]: +

=1

= (1.4 )/( Where the

and

+ 1.4 )

(1) (2)

represent volume fractions of martensite and austenite, respectively;

represents the average of diffraction peak intensities of 200γ, 220γ and 311γ; and

represents

the average of diffraction peak intensities of 200α and 211α. The Modified Williamson–Hall plots [23, 24] were used to evaluate the dislocation density and dislocation character in the RA phase of the specimens. To determine the full width at half 5

maximum (FWHM) and the diffraction angle (θ) of each diffraction peak, the diffraction profiles of 111γ, 200γ, 220γ and 311γ were fitted by Lorentz function. The value of FWHM obtained from each peak is substituted into the following Williamson-Hall equation [23, 24]: ∆ where ∆

= 2cos (∆ )/λ and

.

≌

+

( )

/

= 2sin /λ. Here, ∆

/

(3)

and λ represent the FWHM and

wavelength of X-ray, respectively. The value of λ for Cu radiation was 0.15405 nm. D, b and ρ indicates the average crystallite size, magnitude of the Burgers vector (for γ-Fe b = 2.54×10−10 m) and the dislocation density, respectively. M is a constant depending on both the effective outer cut-off radius of dislocations and the dislocation density [25]. In the case of texture free cubic material, the dislocation contrast factor ( ) for the specific (hkl) reflection can be given by [26]: = & = where

#

#

(1 − %& )

#' (' )#' * ' )(' * ' (#' )* ' )(' )

(4) (5)

is a constant corresponding to the elastic constant of the material. q is a parameter

that depends on the edge or screw character of the dislocation. By combining Eq. (3) and Eq. (4) and substituting

#

, the value of q can be obtained from Eq. (6). (∆+,-)' +'

where / =

.

and . =

( )

/

≌.

#

(1 − %& )

(6)

. The experimental value of α should be determined by

imposing a linear relationship between the left hand term and & . Then, the measured value of q is given by the coefficients of & in the linear function. After obtaining the value of q, the fraction of screw or edge dislocations can be determined by [25]:

6

0 1231 =

:; 456789 ,4 :; :; 456789 ,4<8=8

= 1 − 0 >?@1A

(7)

Where 0 1231 and 0 >?@1A are the fractions of edge and screw dislocations, respectively. Here, %BC# (D=edge or screw) is the theoretical value of q for the pure edge or screw dislocations. The details about the calculation method for %BC# can be found in reference [27]. 2.4 Mechanical properties The Vickers hardness of the tempered specimens was measured by using the HV-1000A Vickers hardness tester and each error bar represents the standard deviation of eight indents. The tensile testing was performed on an electro-hydraulic servo universal testing machine (UTM4503) at a strain rate of ~10-3 s-1. The specimens for tensile testing were prepared by wire-cutting from the tempered specimens as shown in Figure 1.

Figure 1 Sketch of tensile specimen 3. Results 3.1 Microstructure investigations The typical SEM micrographs of the specimens tempered under different conditions are shown in Figure 2. From Figure 2(a), after tempering at 300 °C for 60 min without high magnetic field, the microstructure mainly consists of martensite (M), a substantial amount of blocky RA and a few undissolved carbides (UC). Meanwhile, no obvious carbide precipitation can be observed. However, when high magnetic field is applied, the amount of blocky RA

7

decreases and a few strip-like precipitates are distributed within the martensite matrix (Figure 2(b)). With an increase in tempering temperature up to 380 °C, the strip-like carbides distinctly precipitate from the matrix in the specimens tempered without high magnetic field, as can be seen from Figure 3(c). After the application of high magnetic field (Figure 3(d)), the strip-like precipitates significantly increases compared with those tempered without high magnetic field. Moreover, the RA has been apparently decomposed with only a few blocky RA remaining in the matrix.

Figure 2 SEM micrographs of the specimens tempered at 300 °C for 60 min (a) without and (b) with high magnetic field, and at 380 °C for 60 min (c) without and (d) with high magnetic field. Figure 3(a) exhibits the XRD patterns of specimens tempered under different conditions, where it can be seen that the diffraction peaks of austenite nearly disappear at tempering temperature of 530 °C, regardless of whether high magnetic field is applied. In order to

8

quantitatively investigate the effect of high magnetic field on the decomposition of RA during tempering, the volume fractions of RA for the specimens tempered without and with high magnetic field were measured as shown in Figure 3(b). For the specimens tempered at 300 °C without high magnetic field, the volume fraction of RA remains unchanged compared with that of as-quenched specimens (26.5 %), demonstrating that the RA will not decompose at 300 °C. However, when the tempering temperature increases to 380 °C, the volume fraction of RA shows a reduction in volume fraction without applied high magnetic field (such as, from 26.3 to 21.0 % at holding time of 10 min; from 26.0 to 18.9 % at holding time of 60 min). When the tempering temperature further increases to 530 °C, the volume fractions of RA significantly decreases to 4.2 %, indicating that RA has been almost decomposed completely after the tempering process at 530 °C. However, once the high magnetic field is applied, the volume fraction of RA decreases from ~26 % to ~20 % at the tempering temperature of 300 °C, proving that the decomposition of RA starts from the earlier stage under the high magnetic field. At tempering temperatures of 380 °C and 530 °C, the volume fraction of RA also slightly decreases after applying high magnetic field. Therefore, the high magnetic field can facilitate the decomposition of RA during the tempering process of M50 bearing steel. Furthermore, the dislocation density and dislocation character of RA for the specimens tempered at different conditions, as determined from XRD Rietveld analysis (such as in Figure 3(c)), are shown in Figure 3(d). It can be seen that the dislocation densities of RA in the specimens tempered with high magnetic field are significantly increased, as compared with those in the specimens tempered without high magnetic field. Especially for the specimens tempered at 380 ℃ for 60 min, the dislocation density of RA increases more than doubled from 8.5 × 1015 to 20 × 1015 /m2 after applying high magnetic field. Additionally, it should be noted that the 9

proportion of edge dislocations in RA obviously increases when the high magnetic field is applied.

Figure 3 (a) XRD patterns of specimens tempered under different conditions; (b) The volume fraction of RA under different tempering conditions; (c) XRD pattern and Lorentz fitted result of the specimen tempered at 300 ℃ for 10 min without magnetic field; (d) Dislocation density and dislocation character of RA for the specimens tempered under different conditions. As shown in Figure 4(a), the dilatometry curves and derivative of relative length change scans of the as-quenched specimens were recorded during the non-isothermal heating. It is revealed that the tempering process of M50 bearing steel can be divided into the following sequence of processes. (I) The first stage, referred to as the “pre-precipitation process” and

10

occurring below 100 °C, consists of the segregation of carbon atoms to lattice defects and the formation of clusters of carbon atoms in the matrix. (II) The second stage, associated with the precipitation of transition carbides, occurs in the temperature range of about 80 ~ 200 °C. (III) The third stage, taking place at the temperature above 250 °C, involves the conversion of the transition carbide into cementite and accompanied with a volume contraction. (IV) The fourth stage, occurring at a temperature higher than the low-alloyed steel, relates to the decomposition of retained austenite along with a volume expansion. (V) The fifth stage, assigned to the dissolution of cementite particles and the formation of more stable alloy carbides, takes place at the temperatures above 450 °C.

Figure 4 (a) Dilatometry curves and derivative of relative length change of as-quenched specimens during non-isothermal heating with a rate of 5 °C /min. (b) Derivative of relative length change of the tempered specimens during non-isothermal reheating with a rate of 5 °C /min. Figure 4(b) shows the derivative of relative length change of the tempered specimens during the non-isothermal reheating with a heating rate of 5 °C /min. It can be seen that the pre-precipitation process (stage I) still occurs during the reheating process. After this pre-precipitation process, the derivative of relative length change will approach a constant value 11

(1.51 × 10-5), which is close to the coefficient of thermal expansion (CTE) of martensite (α-Fe) [28]. The results verify that no obvious transformation occurs during the “platform” stage. Meanwhile, a volume contraction, which relates to the conversion of the transition carbide into cementite (stage III), is observed after the “platform” stage. When the tempering temperature increases from 300 °C to 380 °C, the peak of volume contraction almost disappears, indicating that the precipitate of cementite (Stage III) have completed after tempering at 380 ℃. In addition, due to the applied high magnetic field, the peak intensity of volume contraction decreases and the start transformation temperature of cementite precipitation shifts towards higher (as arrowed in Figure 4(b)). This indicates the application of high magnetic field can accelerate the whole tempering process. In other word, the high magnetic field has an effect equivalent to the increase of tempering temperature, which is well consistent with the observation in the literature [29]. During the tempering process, the reduction of interstitial carbon content in the martensite matrix should be mainly caused by the carbide precipitation behaviors [30]. It is known that the diffraction peak information of martensite is closely related to the interstitial carbon content in martensite [31, 32]. To study the effect of high magnetic field on the precipitation behaviors, the magnified XRD profiles of the specimens tempered at 300 °C and 380 °C are fitted by Lorentz function, as shown in Figure 5. It can be seen that the diffraction peaks of martensite (110α, 200α and 211α) shift to higher angles for the specimens tempered under high magnetic field, reflecting that less interstitial carbon atoms exist in the martensite matrix after the application of high magnetic field. The decrease in interstitial carbon atoms could be a consequence of the increase in carbon consumption of cementite precipitate under high magnetic field.

12

Figure 5 Magnified XRD profiles of the specimens tempered at 300 °C (a) 110α, (b) 200α, (c) 211α, and 380 °C (d) 110α, (e) 200α, (f) 211α. To further investigate how the high magnetic field affects the carbide precipitation behaviors, the magnetic hysteresis loops of the carbide powder extracted from the tempered specimens were measured, as shown in Figure 6. For the specimens tempered at 300 °C for 60 min, the magnetization of extracted carbides at a 30 kOe magnetic field increases from 0.17 to 0.84 emu/g after the application of high magnetic field (Figure 6(a)). At higher tempering temperature of 380 °C, the magnetization of extracted carbides is saturated at 30 kOe magnetic field (Figure 6(b)), indicating that the magnetization of precipitated carbides increases obviously. Meanwhile, the magnetization of extracted carbides powder at a 30 kOe magnetic field increases from 1.13 to 5.91 emu/g after the application of high magnetic field. When the specimens are tempered at 530 °C for 120 min, the magnetization of carbides tempered with high magnetic field also increases compared with those tempered without high magnetic field (Figure6(c)).

13

Figure 6 Magnetic hysteresis loops of the carbides powder extracted from the specimens tempered (a) at 300 °C for 60 min, (b) at 380 °C for 60 min and (c) at 530 °C for 120 min As illustrated in Table 2, the Fe and Cr elements of the matrix for the specimens tempered without and with magnetic field were measured based on the solution collected after electrolytic extraction. For the specimens tempered at 300 °C for 60 min, the Fe content in the matrix decreases from 92.06 to 91.20 % while the Cr content in the matrix increases from 4.03 to 4.62 % after applying high magnetic field. When the tempering temperature increases to 380 °C, the Fe content in the matrix of specimens tempered under high magnetic field are also lower than those without magnetic field. For the specimens tempered at 530 °C for 120 min, the Fe content in the matrix obviously increases compared with the specimens tempered at 300 or 380 °C, which may be caused by the increased consumption of alloy elements for the formation of strong alloy 14

carbides. However, due to the application of high magnetic field, the Fe content in the matrix decreases from 94.05 to 93.18 %. Since less Fe content exists in the matrix after tempering with high magnetic field, it can be inferred that the high magnetic field can promote the precipitation of carbides with higher Fe content during tempering. Table 2 The chemical elements of the matrix for the specimens under different tempering conditions Treatment 300 °C - 60 min

380 °C - 60 min

530 °C - 120 min

Magnetic field

Fe (%)

Cr (%)

0T

92.06

4.03

12 T

91.20

4.62

0T

92.13

3.97

12 T

91.61

4.25

0T

94.05

2.56

12 T

93.18

2.93

Figure 7 exhibits the TEM micrographs of specimens after tempering at 300 °C for 60min without and with high magnetic field, where it can be seen that a considerable amount of RA remains undecomposed in the specimens. As shown in Figure 7(a), many ultrafine precipitates distribute within the martensite along the same direction. The martensite martrix is determined to be supersaturated twin martensite (TM) in Figure 7(c). The precipitates, which are identified as ε-carbides (ε-Fe2C) by selected area electron diffraction (SAED) pattern inserted in Figure 7(d), are short rod-lie shape approximately with a length of 150 nm and a width of 20 nm. When high magnetic field is applied, except for the ultrafine ε-Fe2C, some stripe-like carbides with a length of 500 nm are presented within the lath martensite (LM) as shown in Figure 7(b) and (f). By combining SAED analysis inserted in Figure 7(f), the strip-like carbides are identified as cementite (θ-Fe3C), which demonstrates that some ε-carbides have transformed to θ-Fe3C under the action of high magnetic field. As the tempering temperature increases to 380 °C, the massive 15

precipitation of strip-like θ-Fe3C can be clearly observed in Figure 8, regardless of whether high magnetic field is applied. In addition, for the specimens tempered without high magnetic field, the blocky RA identified by SAED pattern locates between the lath martensite (Figure 8(a)). On the contrary, in the specimens tempered with high magnetic field, blocky RA almost disappeared in the microstructure with only a small amount of martensite/austenite (M/A) island existing (Figure 8(b)). With a further increase in the tempering temperature to 530 °C, lath martensite is still dominant in the microstructure (Figure 9). For the specimens tempered without high magnetic field, the dislocation density in lath martensite has been remarkably recovered as illustrated by arrow in Figure 9(a), while there are still extensive dislocations entangled within the lath martensite when the high magnetic field is applied (Figure 9(b)).

Figure 7 TEM micrographs of the microstructures after tempering at 300 °C for 60 min (a) without and (b) with high magnetic field; Dark field images of selected area: (c) A, (d) B, (e) C

16

and (f) D; the insert in (c) is SAED pattern of the twin martensite; the insert in (d) is SAED pattern of the ε-Fe2C; the insert in (e) is SAED pattern of the lath martensite and ε-Fe2C; the insert in (f) is SAED pattern of the θ-Fe3C.

Figure 8 TEM micrographs of the microstructures after tempering at 380 °C for 60min (a) without and (b) with high magnetic field. The insert in (a) is SAED pattern of the martensite and retained austenite; the insert in (b) is the SAED pattern of the martensite and cementite.

Figure 9 TEM micrographs of the microstructures after tempering at 530 ℃ for 120min (a) without and (b) with high magnetic field.

17

3.2 Mechanical properties Figure 10 shows the mechanical properties of the specimens tempered without and with high magnetic field. As illustrated in Figure 10(a), when the specimens are tempered at 300 °C for 60 min, the average Vickers hardness slight increases from 718 to 739 HV after the application of high magnetic field. However, the Vickers hardness for the specimens tempered at 380 °C with high magnetic field is lower than that of without high magnetic field. With a further increase of tempering temperature to 530 °C, the Vickers hardness increases from 764 to 817 HV when the high magnetic field is applied. Figure 10(b) presents the ultimate tensile strength (UTS) varying with different tempering conditions. It can be seen that the UTS decreases during the moderate tempering after the application of high magnetic field (such as, from 1972 to 1851 MPa at 300 °C, and from 1908 to 1783 MPa at 380 °C). However, when the specimens are tempered at 530 °C for 120 min, the UTS increases from 2127 to 2260 MPa after applying high magnetic field. Moreover, comparing the Vickers hardness and UTS at different tempering temperatures, it can be found that the Vickers hardness and UTS first decrease and then increase as the tempering temperature increases from 300 to 530 °C. Figure 10(c) exhibits the results of elongation (EL) for the specimens tempered without and with high magnetic field. It can be found that, due to the application of high magnetic field, the EL decreases obviously at the tempering temperature of 300 and 380 °C. With the temperature further increasing to 530 °C, the presence of high magnetic field during tempering has little effect on the results of EL.

18

Figure 10 The mechanical properties of the specimens tempered without and with high magnetic field: (a) Vickers hardness (b) Ultimate tensile strength (c) Elongation. The tensile fracture morphologies of the specimens tempered at different conditions are shown in Figure 11. After tempering at 300 °C for 60 min, the ductile dimple-like fracture is observed whether the high magnetic field is applied or not (Figure 11(a) and (d)), which is in good agreement with the high elongation at a tempering temperature of 300 °C (Figure 10(c)). When the tempering temperature increases to 380 °C, the fracture surface shows predominant cleavage river pattern with a small quantity of dimples in the core (Figure 11(b)). It indicates a transition of ductile fraction to brittle fraction with the tempering temperature increases from 300 to 380 °C. As for the specimens tempered under a high magnetic field, intergranular fracture can be observed (Figure 11(e)), implying a worse ductile property after the application of high magnetic field at the tempering temperature of 380 °C. When the specimens are subjected to a

19

tempering process at 530 °C, the quasi-cleavage fracture is dominantly observed with few shallow dimples as shown in Figure 11(c) and (f). Therefore, the ductility is slightly improved when the tempering temperature increases from 380 to 530 °C.

Figure 11 Tensile fracture morphologies of the specimens tempered at (a) 300 °C for 60 min, (b) 380 °C for 60 min, (c) 530 °C for 120 min without high magnetic field, and (d) 300°C for 60 min, (e) 380 °C for 60 min, (f) 530 °C for 120 min with high magnetic field. 4. Discussion 4.1 RA decomposition Several researchers [33, 34] have conducted investigations on the effects of high magnetic field on the pre-eutectoid ferrite transformations during an isothermal cooling process in steels. They found that the magnetization of the material coupled with the applied field strength alters the free energy and phase stabilities of austenite, due to the substantial different in magnetic susceptibilities of ferrite and austenite. The decomposition of RA is known as a process of transformation from parent austenite to produced ferrite. During the tempering, RA is a paramagnetic phase while the produced ferrite is a ferromagnetic phase. Thus, there will be a 20

huge difference in magnetic susceptibilities between the two phases under the high magnetic field, which will affect the Gibbs free energy of RA decomposition. The difference in the corresponding Gibbs free energy resulted from the magnetic field can be expressed as Eq. (8) [33]: ∆GF = −(G

FK

N

F IIIIJ H ∙ M IIJ - − G IIIIJ H ∙ M IIJ O )

(8)

Where, the superscripts α and γ denote ferrite and austenite, M and B0 refer to the magnetization and induction of the applied magnetic field, respectively. As the magnetization of ferrite is much higher than that of austenite, the difference in the Gibbs free energy ∆GF is negative. Consequently, the application of high magnetic field will result in a decrease of the energy barrier of transformation from RA to ferrite. This corresponds a promotion of decomposition of RA induced by the high magnetic field during tempering, as illustrated in Figure 3(b). In this research, a new mechanism to explain why the decomposition of RA is facilitated by the application of high magnetic field is proposed. Magnetic field-induced deformation (magnetoplasticity), as one of important magneto-mechanical behaviors of ferromagnetic material, is generally realized by the field-induced displacement of ordinary dislocation or twin dislocation [35, 36]. Under the applied high magnetic field, the magnetic field-induced deformation of martensite will inevitably affect the neighboring austenite. The inhomogeneous plasticity generated from the deformed martensite may lead to the initial and motion of dislocation in ductile austenite phase, thereby increasing the dislocation density in RA (Figure 3(d)). Moreover, it should be noted that edge type dislocation dominates plasticity at ambient temperature and 300 °C, while screw type dislocation dominates at 600 °C [37]. This could explain why the proportion of edge dislocations in the RA increases after the application of high magnetic field, in the context of a moderate temperature tempering (at 300 or 380 °C).

21

As well known, the degree of RA decomposition is directly determined by the thermal stability of RA during the tempering process. In recent years, some researchers [38, 39] have studied the relationship between the thermal stability of RA and dislocation density in RA. They point out that defects such as dislocations and shear bands in austenite are beneficial to the nucleation and formation of martensite, thus the increase of dislocation density in RA correlates with a decrease of austenite stability. Accordingly, the promoted decomposition of RA under the high magnetic field could be a consequence of the increasing dislocation density in RA due to the magnetic field-induced deformation of martensite. 4.2 Carbide precipitation The above microstructural observation and dilatometry analysis have shown that the precipitate of cementite, which corresponds to the stage III above 250 °C, is promoted by the applied high magnetic field after tempering at 300 °C and 380 °C. Based on the classical nucleation theory [40], the homogeneous nucleation rate per volume can be described as Eq. (9): P = P exp T−

U

VW

X exp (−

∆Y ∗ VW

)

(9)

where P is a constant, [, R and T represent the diffusion activation energy, universal gas constant and the absolute temperature, respectively. The nucleation barrier ∆\ ∗ reads [41]: ∆\ ∗ = where

]^

∆Y_'

(10)

is a constant, ` is the interfacial energy between the carbide and the matrix, and ∆\a

represents the Gibbs volume free energy difference between the two phases (ferrite and cementite) or driving force of the transformation. The negative term ∆\a can be obtained as [42]: ∆\a = \?1b1cCBC1 − \d1@@BC1 < 0

22

(11)

As the Curie temperature of cementite is 225 °C [43], the magnetism of cementite will change from ferromagnetism to paramagnetism at the tempering temperature of 300 °C or 380 °C. When the high magnetic field is applied, it will result in a larger decrease of \?1b1cCBC1 than that of \d1@@BC1 due to the transformation of ferromagnetism to paramagnetism of ferrite. As a result, an increase of the absolute value of ∆\a can be expected under the high magnetic field, which will give rise to the reduction in the total nucleation barrier ∆\ ∗ in Eq. (10). Finally, the lower nucleation barrier leads to higher nucleation rate P as illustrated in Eq. (9). In present work, the dilatometry peak of cementite precipitation (stage ℃) almost disappears after tempering at 380 °C (Figure 4b), demonstrating that cementite has completely precipitated at the given tempering temperature. However, it is worth noting that the magnetizations of extracted carbide for the specimens tempered without and with high magnetic field are significantly different, as shown in Figure 6(b). It has been reported [44, 45] that θ-Fe3C have the highest magnetization at ambient temperature, as compared with other alloy carbides. After a tempering process of 380 °C, the carbide powder obtained by extraction is mainly composed of cementite and undissolved carbides. The undissolved carbides of M50 bearing steel after a high temperature austenitization have been identified as V-rich MC and Mo-rich M2C [46], which has a much lower magnetization than cementite. Therefore, the improvement of magnetic properties of extracted carbide may be attributed to the higher proportion of cementite under high magnetic field, which is consistent with the observations in Figure 4 and Figure 5. In conclusion, the total amount of precipitated cementite is significantly increased by the high magnetic field. When the tempering temperature reaches 530 °C, the strong alloy carbides (V-rich MC Mo-rich M2C, Cr-rich M7C3 and M23C6) form in M50 bearing steel [47]. Recently, many researches have proven [12, 48-50] that the amount of substitutional solute atoms in alloy

23

carbides has been found to greatly affect the magnetic properties of carbides, thus determining their precipitation behaviors. For example, Hou et. al [48] reported that the concentration of substitutional solute atoms Fe in the precipitated Cr-rich carbides was greatly influenced by the high magnetic field. And they stated that higher Fe atom content in the Cr-rich carbide would cause a remarkable increase in the magnetic moment, which resulted in a reduction in the magnetic Gibbs free energy of this carbide. Zhou et. al [12] hold that the atomic ratio of Fe/Mo was greatly increased by the application of high magnetic field for the molybdenum carbides, thus leading to an improvement of precipitation of the Mo-containing with a higher substitutional solute atoms Fe. It thus can be expected that the magnetization of carbides is mainly contributed by the Fe content in carbides and more substitutional solute atoms Fe in carbides will increase the magnetization of carbides. Other interstitial elements such as C and Cr, however, could decrease the magnetic properties of carbides with an addition of content [49, 50]. Under the high magnetic field, because of the decrease of Fe content and increase of Cr content in the matrix after tempering (Table 2), it is inferred that alloy carbides with higher Fe content will precipitate preferentially during tempering at 530 °C in M50 bearing steel. 4.3 Mechanical properties In order to investigate the effect of high magnetic field on the mechanical properties of M50 bearing steel after different tempering process, the results of hardness, tensile experiments and tensile fracture were analyzed for the specimens tempered with and without magnetic field. When the specimens are tempered at 300 °C for 60 min, the Vickers hardness of specimens slightly increases after applying high magnetic field (Figure 10(a)). Whereas, the tensile strength and UTS of the specimens tempered with high magnetic field show decrease trends compared with those without magnetic field (Figure 10(b) and (c)). The differences in mechanical

24

properties induced by high magnetic field should be mainly related to three aspects: (a) decomposition of soft phase RA, (b) conversion of the transition ε-carbide into strip-like cementite, and (c) the solid solution strengthening effect of martensite. After applying the high magnetic field, the significant decrease in volume fraction of RA are in favor of enhancing the hardness, whereas the increased precipitate of cementite and the decreased solid solution carbon are unfavorable for the hardness. It is thus found that the increase of Vickers hardness, which corresponds to a decrease in the EL, should be mainly attributed to the decomposition of RA when the specimens are tempered at 300 °C. It has been proved that the strength decreases obviously when the ε-carbide is transformed into cementite [51]. Therefore, the accelerated formation of cementite induced by high magnetic field plays a dominant role in decreasing the UTS during tempering at 300 °C. With an increase of tempering temperature to 380 °C, the Vickers hardness, UTS and EL all decrease compared with those tempered at 300 °C. Meanwhile, the application of high magnetic field aggravates the deterioration of mechanical properties. The M50 bearing steel, as a high strength martensitic steel, is susceptible to embrittlement after tempering at certain temperatures. This phenomenon is generally called as tempered martensite embrittlement (TME) [52], which is demonstrated to be closely related to the formation of platelet cementite (which replaced ε-carbide) [53] and decomposition of RA [54, 55]. Since the precipitation of cementite and decomposition of RA are enhanced by the application of high magnetic field, it is not difficult to understand that the EL of specimen is the lowest after tempering at 380 °C with the high magnetic field. Furthermore, the solid solute carbon in the martensite matrix, as crucial factors affecting the hardness and UTS [56], decreases due to the promoted precipitation of cementite by high magnetic field. Accordingly, the decreases in hardness and UTS after the application of high

25

magnetic field should be mainly caused by the reduction in solution strengthening of martensite phase. When the specimens are tempered at 530 °C for 120 min, significant increases in Vickers hardness and UTS can be found compared with those tempered at 300 °C or 380 °C. This reveals that the obvious secondary hardening occurs during a tempering process of 530 °C. Meanwhile, it indicates that the Vickers hardness and UTS increase after applying high magnetic field. Figure 9 has shown that the dislocation recovery in the martensite is retarded by high magnetic field. In fact, the similar delaying effect caused by the high magnetic field on recovery process has been observed in the cold rolled steel [57, 58] and as-quenched martensitic steel [16]. It is stated the domain walls may act as barriers to grain or sub-grain boundary migration, thus delaying the processes of dislocation recovery. Consequently, the increase of Vickers hardness and UTS should be mainly attributed to the higher dislocation strengthening after applying high magnetic field. The tensile fracture morphologies of the specimens tempered without and with high magnetic field are presented in Figure 11. With the increase of tempering temperature from 300 °C to 380 °C, a transition of ductile fraction to brittle fraction can be clearly observed, which could be attributed to the TME effect as discussed before. It is noted that the tensile fracture morphology changes from cleavage river fracture pattern to intergranular fracture pattern after the application of high magnetic field (Figure 11(b) and (e)). It means that the high magnetic field significantly intensifies the TME effect at a tempering temperature of 380 °C. With a further increase in tempering temperature to 530 °C, the quasi-cleavage fracture is dominantly observed with few shallow dimples in the tensile fracture (Figure 11(c) and (f)). Therefore, the tempering process of 530 °C can avoid the TME effect in M50 bearing steel, but the ductility is

26

still poor because the volume fraction of RA is only about 4%. 5. Conclusion In this work, the microstructure evolution during tempering of M50 bearing steel have been investigated under the high magnetic field. The mechanical properties of specimens tempered without and with high magnetic field have been discussed in detail. The following conclusion can be obtained. (a) The decomposition of RA during a moderate tempering (at 300 or 380 °C) is promoted by high magnetic field. The thermodynamic analysis reveals that the energy barrier of transformation from RA to ferrite decreases to accelerate the decomposition of RA due to the difference in magnetic susceptibilities between austenite and ferrite. Meanwhile, because of the magnetic field-induced deformation of martensite, an obvious increase in the dislocation density of RA is observed after applying high magnetic field. This increased dislocation density provides more nucleation site for the transformation from RA to ferrite, thus contributing to the decomposition of RA. (b) During a moderate tempering, the total amount of cementite precipitated is significantly increased by the high magnetic field due to the higher kinetic nucleation rate. The contribution of high magnetic field to cementite precipitation is verified by the results of carbon solid solute in martensite and magnetization of precipitated carbides. When the specimens are tempered at 530 °C for 120 min, the Fe content in the matrix decreases while the Cr content increases after applying high magnetic field. The magnetic properties of carbides are greatly influenced by the substitutional solute atoms in alloy carbides, leading to the preferential precipitate of the alloy carbides with higher Fe content.

27

(c) During tempering at 300 °C, the promoted RA decomposition by high magnetic field results in a slight increase in Vickers hardness. Additionally, the increased formation of cementite induced by high magnetic field plays a dominant role in the decrease of UTS from 1972 to 1851 MPa. With an increase in the tempering temperature to 380 °C, the TME effect is significantly intensified due to the accelerated RA decomposition and improved cementite precipitation by the high magnetic field. When the specimens are tempered at 530 °C, as the dislocation recovery is retarded by high magnetic field, the Vickers hardness increases from 764 to 817 HV and the UTS increases from 2127 to 2260 MPa after applying high magnetic field. This finding thus suggests the potential of high magnetic field (12T) in optimizing the properties of M50 bearing steel during tempering of 530 °C.

Acknowledgements This work was financially supported by National Natural Science Foundation of China (51575414, 51875426), Innovative Research Team Development Program of Ministry of Education of China (No. IRT13087) and 111 Project (B17034).

Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

Reference [1] Y. Zhang, C. Esling, Phase Transformations in Steels: The Effect of a Magnetic Field on

28

Phase Transformations in Steels, Woodhead publishing, 2012, pp. 555-580. [2] G. M. Ludtka, R. A. Jaramillo, R. A. Kisner, D. M. Nicholson, J. B. Wilgen, G. Mackiewicz-Ludtka, P. N. Kalu, In Situ Evidence of Enhanced Transformation Kinetic in a Medium Carbon Steel due to a High Magnetic field. Scr. Mater. 51 (2004) 171-174. [3] S. Y. Yu, S. S. Yan, S. S. Kang, X. D. Tang, J. F. Qian, J. L. Chen, G. H. Wu, Magnetic Field-induced Martensite–austenite Transformation in Fe-substituted NiMnGa Ribbons, Scr. Mater. 65 (2011) 9-12. [4] J. Y. Choi, T. Fukuda, T. Kakeshita, Effect of Magnetic Field on Isothermal Martensitic Transformation in a Sensitized SUS304 Austenitic Stainless Steel, J. Alloy. Comp. 577S (2013) S605–S608. [5] D. S. Martin, N. H. V. Dijk, E. J. Melero, E. Kampert, U. Zeitler, S. V. D. Zwaag, Real-time Martensitic Transformation Kinetics in Maraging Steel under High Magnetic Fields, Mater. Sci. Eng. A. 527 (2010) 5241-5245. [6] R. A. Jaramillo, S. S. Babu, G. M. ludtka, R. A. Kisner, J. B. Wo;gen, G. M. Ludtka, D. M. Nicholson, S. M. Kelly, M, Murugananth, H. K. D. H. Bhadeshia, Effect of 30 T Magnetic Field on Transformations in a Novel Bainitic Steel, Scr. Mater. 52 (2005) 461-466. [7] B. Q. Dong, T. P. Hou, K. M. Wu, Z. Q. You, Z. H. Li, G. H. Zhang, H. F. Lin, Low-temperature Nanostructured Bainite Transformation: The Effect of Magnetic Field, Mater. Lett. 240 (2019) 66-68. [8] T. Garcin, S. Rivoirard, C. Elhoyhen, E. Beaugon, Experimental Evidence and Thermodynamics Analysis of High Magnetic Field Effects on the Austenite to Ferrite Transformation Temperature in Fe–C–Mn Alloys, Acta. Mater. 58 (2010) 2026-2032. [9] Y. D. Zhang, C. S. He, X. Zhao, L. Zuo, C. Esling, J. C. He, New Microstructural Features

29

Occurring during Transformation from Austenite to Ferrite under the Kinetic Influence of Magnetic Field in a Medium Carbon Steel, J. Magn. Magn. Mater. 284 (2004) 287-293. [10] Y. D. Zhang, C. Esling, M. L. Gong, G. Vincent, X. Zhao, L. Zuo, Microstructural Features Induced by a High Magnetic Field in a Hypereutectoid Steel during Austenitic Decomposition, Scr. Mater. 54 (2006) 1897-1900. [11] J. J. Li, W. Liu, Effect of High Magnetic Field on Isothermal Pearlite Transformation and Microstructure in a Hypereutectoid Steel. J. Magn. Magn. Mater. 362 (2014) 159-164. [12] J. Y. Song, Y. D. Zhang, X. Zhao, L. Zuo, Effects of high magnetic field strength and direction on pearlite formation in Fe–0.12%C steel. J. Mater. Sci. 43 (2008) 6105-6108. [13] Z. N. Zhou, K. M. Wu, Molybdenum Carbide Precipitation in an Fe–C–Mo Alloy under a High Magnetic Field, Scr. Mater. 61 (2009) 670-673. [14] T. P. Hou, K. M. Wu, Alloy Carbide Precipitation in Tempered 2.25 Cr–Mo Steel under High Magnetic Field, Acta. Mater. 61 (2013) 2016-2024. [15] G. H. Wu, T. P. Hou, K. M. Wu, L. Chen, Influence of High Magnetic Field on Carbides and the Dislocation Density during Tempering of High Chromium-containing Steel, J. Magn. Magn. Mater. 479 (2019) 43-49. [16] Y. D. Zhang, N. Gey, C. S. He, X. Zhao, L. Zuo, C. Esling. High Temperature Tempering Behaviors in a Structural Steel under High Magnetic Field, Acta. Mater. 52 (2004) 3467-3474. [17] Y. V. Kaletina, E. A. Fokina, Effect of magnetic field on the amount of retained austenite in hardened steels, Met. Sci. Heat. Treat. 50 (2008) 482-488. [18] E. A. Fokina, V. N. Olesov, L. V. Smirnov, Effect of Magnetic Field on Transformation of Retained Austenite in Steels, Met. Sci. Heat. Treat. 18 (1976) 109-110.

30

[19] J. I. Pe´rez-Landaza´bal, V. Recarte, V. Sa´nchez-Alarcos, Go´mez-Polo, S. Kustov, E. Cesari, Magnetic Field Induced Martensitic Transformation Linked to the Arrested Austenite in a Ni-Mn-In-Co Shape Memory Alloy, J. Appl. Phys. 109 (2011) 1-5. [20] A. Dodd, J. Kinder, B. Torp, B.R. Nielsen, C.M. Rangel, M.F.da Silvad, The Effect of Ion Implantation on the Fatigue Life and Corrosion Resistance of M50 Steel Bearings, Surf. Coat. Technol. 74-75 (1995) 754-759. [21] H.K.D.H. Bhadeshia, Steels for Bearings, Prog. Mater. Sci. 57 (2012) 268-435. [22] Y. K. Lee, H. C. Shin, Y. C. Jang, S. H. Kim, C. S. Choi, Effect of Isothermal Transformation Temperature on Amount of Retained Austenite and Its Thermal Stability in a Bainitic Fe–3%Si–0.45%C–X Steel, Scr. Mater. 47 (2002) 805-809. [23] T. Shintani, Y. Murata, Evaluation of the Dislocation Density and Dislocation Character in Cold Rolled Type 304 Steel Determined by Profile Analysis of X-ray Diffraction, Acta. Mater. 59 (2011) 4314-4322. [24] T. Ungár, Dislocation densities, arrangements and character from X-ray diffraction experiments, Mater. Sci. Eng. A. 309-310 (2001) 14-22 [25] F. Hajyakbary, J. Sietsma, A. J. Böttger, M. J. Santofimia, An improved X-ray diffraction analysis method to characterize dislocation density in lath martensitic structures, Mater. Sci. Eng. A. 639 (2015) 208-218. [26] F. Yin, T. Hanamura, O. Umezawa, K. Nagai, Phosphorus-induced dislocation structure variation in the warm-rolled ultrafine-grained low-carbon steels, Mater. Sci. Eng. A. 354 (2003) 31-39 [27] T. Ungár, I. Dragomir, Á. Rèvèsz, A. Borbèly, The contrast factors of dislocations in cubic crystals: the dislocation model of strain anisotropy in practice, J. Appl. Cryst. 32 (1999)

31

992-1002. [28] Thermal Expansion characteristics of Fe–9Cr–0.12C–0.56Mn–0.24V–1.38W–0.06Ta (wt.%) Reduced Activation Ferritic–martensitic Steel. J. Nucl. Mater. 459 (2015) 150-158. [29] Y. D. Zhang, N. Bozzolo, C. S. He, X. Y. Zhao, L. Zuo, C. Esling, J. C. He, G. Hishijima, T. Zhang, Influence of High Magnetic Field on the Precipitation in a Medium Carbon Low-Alloyed Steel, Second International Conference on Advanced Structural Steels, (2004) 651-660. [30] J. J. Hoyos, A. A. Ghilarducci, D. Mari, Evaluation of Dislocation Density and Interstitial Carbon Content in Quenched and Tempered Steel by Internal Friction. Mater. Sci. Eng. A. 640 (2015) 460-464. [31] X. Liu, F. Zhong, J. X. Zhang, Lattice-parameter Variation with Carbon Content of Martensite. I. X-ray-diffraction Experimental Study. Phys. Rev. B. 52 (1995) 9970–9978. [32] F. Wang, D. S. Qian, L. Hua, H. J. Mao, Lechun Xie, Voids healing and carbide refinement of cold rolled M50 bearing steel by electropulsing treatment. Sci. Rep-UK. 9 (2019) 11315. [33] Y. D. Zhang, C. S. He, X. Zhao, L. Zuo, C. Esling, Thermodynamic and Kinetic Characteristics of the Austenite-to-ferrite Transformation under High Magnetic Field in Medium Carbon Steel. J. Magn. Magn. Mater. 294 (2005) 267-272. [34] H. D. Joo, S. U. Kim, Y. M. Koo, N. S. Shin. J. K. Choi, An Effect of a Strong Magnetic Field on the Phase Transformation in Plain Carbon Steels. Metall. Mater. Trans. A. 35 (2004) 1663-1668. [35] P. Müllner, V. A. Chernenko, G. Kostorz, A Microscopic Approach to the Magnetic-field-induced Deformation of Martensite (Magnetoplasticity). J. Magn. Magn. Mater. 267 (2003) 325-334.

32

[36] H. Karo, K. Sasaki, Straining of Magnetic Shape Memory Martensite by Uniform Magnetic Field. Scr. Mater. 48 (2003) 31-35. [37] G. M. Zhang, Z. J. Zhou, K. Mo, Y. B. Miao, X. Liu, J. Almer, J. F. Stubbins, The Evolution of Internal Stress and Dislocation During Tensile Deformation in a 9Cr Ferritic/Martensitic (F/M) ODS Steel Investigated by High-energy X-rays. J. Nucl. Mater. 467 (2015) 50-57. [38] J. Hidalgo, K. O. Findley, M. J. Santofimia, Thermal and Mechanical stability of Retained Austenite Surrounded by Martensite with Different Degrees of Tempering. Mater. Sci. Eng. A. 690 (2017) 337-347. [39] X. Qiao, L. Z. Han, W. M. Zhang, J. F. Gu, Thermal Stability of Retained Austenite in HighCarbon Steels during Cryogenic and Tempering Treatments. Isij. Int. 56 (2016) 140-147. [40] H. L. Zhang, X. W. Tang, R. H. Wei, S. J. Zhu, J. Yang, W. H. Song, J. M. Dai, X. B. Zhu, Y. P. Sun, Microstructure refinement and magnetization improvement in CoFe thin films by high magnetic field annealing. J. Alloy. Compd. 729 (2017) 730-734. [41] Y. D. Zhang, C. S. He, X. Zhao, L. Zuo, C. Esling, J. C. He, New microstructural features occurring during transformation from austenite to ferrite under the kinetic influence of magnetic field in a medium carbon steel. J. Magn. Magn. Mater. 284 (2004) 287-293. [42] B. Z. Cui, K. Han, H. Garmestani, J. H. Su, H. J. Schneider-Muntau, J. P. Liu, Enhancement of exchange coupling and hard magnetic properties in nano-composites by magnetic annealing. Acta. Mater. 53 (2005) 4155-4161. [43] D. Chaira, B. K. Mishra, S. Sangal, Magnetic properties of cementite powder produced by reaction milling. J. Alloy. Compd. 474 (2009) 396-400. [44] M. A. Konyaeva, N. I. Medvedeva, Electronic Structure, Magnetic Properties, and Stability of the Binary and Ternary Carbides (Fe, Cr)3C and (Fe, Cr)7C3. Phys. Solid. State. 51 (2009)

33

2084-2089. [45] T. P. Hou, K. M. Wu, W. M. Liu, M. J. Peet, C. N. Hulme-Smith, L. Guo, L. Zhuang. Magnetism and High Magnetic Field-induced Stability of Alloy Carbides in Fe-based Materials. Sci. Rep-UK. 9 (2008) 3049. [46] F. Wang, D. S. Qian, L. Hua, X. H. Lu, The effect of prior cold rolling on the carbide dissolution, precipitation and dry wear behaviors of M50 bearing steel, Tribol. Int. 132 (2019) 253-264. [47] J. E. Bridge, G. N. Maniar, T. V. Philip, Carbides in M-50 high speed steel, Metall. Mater. Trans. B. 2 (1971) 2209-2214. [48] T. P. Hou, K. M. Wu, The effect high magnetic field on metal solute substitution in M23C6 alloy carbide, Scr. Mater. 67 (2012) 609-612. [49] X. B. Wang, K. Zhu, Y. M. Ju, Y. Li, W. Li, J. J. Xu, Y. L. Hou, Iron carbides: Magic materials with magnetic and catalytic properties. J. Magn. Magn. Mater. 489 (2019) 165432. [50] B. Wang, Q. Zhang, Z. F. Zhang, Z. Q. Lv, W. T. Fu, Effect of Cr on electronic and magnetic properties of χ-carbide (Fe, Cr)5C2, J. Magn. Magn. Mater. 392 (2015) 1-5. [51] F. G. Liu, X. Liu, M. H. Song, H. Yang, K. Song, P. F. Guo, W. D. Huang, Effect of tempering temperature on microstructure and mechanical properties of laser solid formed 300M steel, J. Alloy. Compd. 689 (2016) 225-232. [52] H.K.D.H. Bhadeshia, R. Honeycombe, Steels: Microstructure and Properties (Fourth edition), 2017, PP. 237-270. [53] L. J. Klingler, W. J. Barnett, R. P. Frohmberg. A. R. Troiano, The embrittlement of alloy steel at high strength levels, Trans. ASM. 46 (1954) 1557. [54] J. P. Materkowski, G. Krauss, Tempered martensite embrittlement in SAE 4340 steel, Metall.

34

Trans. A. 10A (1979) 1643-1651. [55] G. Krauss, Steels: Processing, Structure, and Performance, ASM International, Materials Park, OH, 2005. [56] Y. P. Zhang, D. P. Zhan, X. W. Qi, Z. H. Jiang, Effect of Tempering Temperature on the Microstructure and Properties of Ultrahigh-strength Stainless Steel, J. Mater. Sci. Technol. 35 (2019) 1240-1249. [57] H. O. Masahashi, V. K. Lindroos, Observations of the Effect of Magnetic-field on the Recrystallization in Ferrite, Scand. J. Metall. 10 (1981) 3-8. [58] C. S. He, Y. D. Zhang, X. Zhao, L. Zuo, J. C. He, K. Watanabe, et al. Effects of a High Magnetic Field on Microstructure and Texture Evolution in a Cold-Rolled Interstitial-Free (IF) Steel Sheet during Annealing, Adv. Eng. Mater. 5 (2003) 579-583.

35

Declaration of Interest Statement

The authors declared that they have no conflicts of interest to this work.