Carbide precipitation during tempering of hybrid steel 60

The effects of carbide precipitation on mechanical performance of Hybrid Steel 60, known as a novel bearing steel, have not been investigated. In this study, the austenite transformation temperatures of Hybrid Steel 60 during heating were revealed by the thermal expansion curve. The temperature and effective activation energy of the second phase precipitation were determined by the differential scanning calorimetry (DSC) curve. Different solid solution structures after austenitization were detected using various cooling rates. The solubility temperature was determined based on hardness and residual austenite content. The carbides precipitated at the peak temperature were qualitatively identified using XRD. It was discovered that the temperature points Ac1 and Ac3 of the steel were 786 °C and 864 °C, respectively. In addition, the effect of solid solution temperature on quenching hardness is minimal, while the cooling rate has a greater impact on hardness, reaching a peak at 5 °C s−1. The primary carbide phase in Hybrid Steel 60 is the M7C3 and VC. When the temperature ranges from 500 °C to 550 °C, M23C6 begins to precipitate. As a result, after tempering at 525 °C, the hardness peak value reached 566 HV.


Introduction
Bearings play a vital role as strategic materials in the national economy and they are crucial components of the manufacturing industry.To meet the demands of increasingly complex transportation systems, there is a growing need for higher-quality bearing steel [1,2].Initially, the high-temperature resistant steels, such as American Iron and Steel Institute (AISI) 52100, which can withstand temperatures of 150 °C, were employed as aviation bearing steels [1].Following World War II, second-generation bearing steels, including M50 and carburized M50NiL (primarily used for bearing steel rings), emerged at a 350 °C usage temperature [3][4][5].In the 1990s, with the development of computational phase diagrams, the latest third-generation bearing steel, CSS-42L, was introduced; this steel boasted more prominent high-temperature hardness values than its predecessors, reaching 453.8 HV after tempering at 580 °C [6,7].In recent years, to accommodate increasingly stringent service conditions, a novel bearing steel with a yield and tensile strength more than double that of traditional bearing steel AISI 52100 at high temperatures has emerged.The steel hardness can reach 566 HV after tempering at 525 °C, making it an ideal choice for various high-stress applications [8].
The second-phase reinforcement constituent of this new bearing steel encompasses not only carbides, but also intermetallic compounds, which distinguish it from conventional bearing steel.The tempering of traditional martensitic steel involves four overlapping stages: carbon segregation, carbide precipitation, residual austenite decomposition, and ferrite recovery.In tempered martensite, second-phase strengthening, which is mainly composed of carbide [9], plays an important role.As the temperature increases, carbides continuously precipitate and grow, gradually decreasing the hardness [10,11].The final strength of hybrid steel is derived from a martensitic microstructure containing a combination of alloy carbides and intermetallic precipitation hardening.The intermetallic precipitate has a body-centered cubic (BCC)/B2 crystal structure and is composed of (Fe, Ni)Al.The carbides present are plate-shaped MC and spherical M 7 C 3 precipitates [8].In standard steel, carbon segregation ordinarily occurs during the initial tempering stage.Carbon atoms are allocated to lowenergy sites, such as vacancies, dislocations, and grain boundaries.Subsequently, transitional carbides, including ε-carbide (Fe 2.4 C), η-carbide and χ-carbide, form at carbon-enriched sites.Exceeding 250 °C results in the formation of θ-carbide (Fe 3 C), which is also known as cementite [12].
For alloy steel, as the tempering temperature increases, alloy carbides, such as M 23 C 6 and M 7 C 3 (M represents metallic components, such as Fe, Cr and Mo, in carbides), and MC emerge [13].The morphology, type, distribution, and quantity of carbides in steel significantly influence its properties.Therefore, examining carbides during the tempering process is vital for attaining superior mechanical properties and offers guidance for material heat treatment [14,15].The aims of the present study are to scrutinize the quenching and tempering behavior of Hybrid Steel 60 at various temperatures, to analyze its carbide transformation following tempering, and to investigate the optimal temperature for heat treatment by assessing the microstructural and hardness attributes of the material.These results should provide basic data for the design of reasonable heat treatment schemes and organizational optimization schemes for Hybrid Steel 60.

Material composition and preparation
Hybrid Steel 60 bearing steel was used as the experimental material, and its chemical composition was measured by a mobile spectrum analyzer, as shown in table 1.
The material was cut into 4 thermal specimens for analysis, each with dimensions of Ø4 × 2 mm, labeled a, b, c, and d.During preprocessing, samples a and c were treated as experimental samples, heated to 920 °C, and quenched.Samples b and d were treated as reference samples and subjected to annealing at 760 °C after quenching at 920 °C.During the experiment, pretreated sample a was placed in a sample crucible, pretreated sample b was placed in a reference crucible, and both samples were heated to 1100 °C at a rate of 5 °C s −1 .Similarly, pretreated sample c was placed in a sample crucible, and pretreated sample d was placed in a reference crucible; both were heated to 1100 °C at a rate of 20 °C s −1 .The heat treatment process is shown in figure 1.The differential scanning calorimetry (DSC) curves of the specimens were obtained using an STA 409PC synchronous thermal analyzer (Netzsch, Selb, Germany) with argon as the protective gas during the test.
The material was cut into several Ø4 × 10 mm samples, heated to 920 °C using a DIL 402 C dilatometer (Netzsch, Selb, Germany) at a rate of 5 °C s −1 , maintained at that temperature for 15 min, and subsequently cooled at a rate of 0.2 °C s −1 to obtain the expansion curve, as shown in figure 2(a) to clarify the Ac1 and Ac3 of the material.
Subsequently, the experimental material samples were quenched in water after they had held for 30 min in a vacuum furnace at 900 °C, 920 °C, 950 °C, 1000 °C, 1050 °C, and 1100 °C, after which the hardening ability and residual austenite content of the samples were analyzed.For tempering, the samples that were air cooled after vacuum quenching after holding at 920 °C for 1 h were held in a vacuum furnace for 3 h at 450 °C, 500 °C, 525 °C, 550 °C, 600 °C, 650 °C, and 760 °C before air cooling, after which the hardness and microstructure were analyzed.

Microscopic characterization
The metallographic sample was prepared using standard grinding and polishing techniques, followed by etching with an etchant consisting of 150 ml of water, 50 ml of hydrochloric acid, and 25 ml of nitric acid.Etching with Kalings Reagent 2 (100 ml absolute alcohol + 100 ml HCl + 5 g CuCl 2 ) when observing residual austenite.Use Leica DMI5000M optical microscope for observation.The microstructures of samples subjected to different heat treatment processes were observed using a SUPRA40-41-90 (Zeiss, Oberkochen, Germany) field emission scanning electron microscope (SEM).The hardness of quenched and tempered samples was measured using a QATM automatic hardness tester with a 1 kg load and a 10 s loading time.Each hardness point represented a mean value of five measurements, and the error is within ± 10 HV.The residual austenite content at different quenching temperatures was measured using an X-350A stress tester (AST, Handan, China).During the test, the high voltage of the x-ray tube was 20 KV, and the current of the X-ray tube was 5 A. When measuring, the 2θ scanning range of martensite (ferrite) is 169°to 142°, the scanning step is 0.2°, and the counting time is 0.5 seconds; The 2θ scanning range of austenite is 134°to 123°, with a scanning step of 0.1°and a counting time   of 1 s.To ensure the accuracy of the results, the samples were polished and polished before testing, and the average value was taken after repeated testing.The electrolyte was prepared using a concentration of 10% HCl and 5% citric acid, and a current density of 0.5 A cm −2 ; the test steel was used as an anode, with a platinum electrode as a cathode.Phase analysis was performed using a Smartlab 9 kW high-resolution x-ray diffractometer (Rigaku, Japan) with a technical index of 40 kv-100 mA, Cu target, 0.02°scanning step, and 0.4°min −1 scanning speed; raw data were compared with the PDF card using Jade.

Structure after quenching
Quenching and tempering are the predominant techniques employed to optimize the thermal processing of a specific steel to enhance its properties.The microstructure of steel after quenching significantly influences its microstructure and performance after tempering.Therefore, before examining a tempering process, it is essential to evaluate the performance of steel after quenching [16].Appropriate quenching requires determining the austenitizing temperature of steel.Based on figure 2(a), the Ac1 and Ac3 temperatures of the sample material are determined to be 786 °C and 864 °C, respectively.Thus, the tested hardening capacity should be greater than the Ac3 temperature.According to figure 2(b), the residual austenite content after quenching increases as the austenitizing temperature increases, indicating that more alloying elements melt into austenite within this temperature range, gradually increasing the austenite stability.Therefore, to examine the hardening capacity, it is necessary to quench a sample following an extended duration of holding the temperature.
Based on the metallographic structure shown in figures 3 and 4, the quenching structure heated from 900 °C to 950 °C consists of martensite, residual austenite and carbides.In addition, the austenite grain size increases as the quenching temperature increases.After quenching at 1000 °C, the metallographic structure changes, as shown in figure 3(d).Based on the hardness characteristics of alloyed steel, the alloy carbides begin to dissolve, and the alloying elements enter the austenite phase.The austenite then cools to form hard, oversaturated martensite.Incorporating alloying elements improves the supercooled austenite stability and prevents carbide precipitation.Therefore, the metallographic structure contains residual austenite, undissolved alloy carbides, and martensite while maintaining the austenite grain morphology [17].
To improve the high-temperature hardness of steel, it is necessary to retain alloy carbides and intermetallic compounds during austenitizing to increase the number of nucleation sites available for carbide precipitation during tempering and thus to obtain a more uniform carbide distribution.Furthermore, it is important to retain some residual austenite to maintain the toughness of steel.Therefore, based on comprehensive consideration, the samples are held at 920 °C for 1 h or 1.5 h, followed by water quenching.The hardness of both samples is 459 HV, indicating that the matrix is fully austenitized after 1 h at 920 °C and that the structure is relatively uniform.The expansion curve shows that the material undergoes martensitic transformation, but its room temperature hardness (figure 2(b)) is much lower than the expected hardness of quenched martensite.This finding indicates that the martensite transformation is not completed under water quenching conditions.This trend occurs because alloying elements (such as Cr and Mo) in the material increase the supercooled austenite stability.A comparison of the expansion curves under different cooling rates (figure 5) reveals that the starting point of martensitic transformation decreases as the cooling rate increases and that the end point of martensitic transformation is pushed below room temperature, resulting in incomplete martensitic transformation [18,19].To ensure the percentage of martensitic transformation and solid solution strengthening effect, air cooling should be used for quenching.

DSC and tempering hardness
The specimens undergo distinct cooling rates followed by a stepwise tempering process ranging from 450 °C to 750 °C. Figure 6(a) shows the hardness results, indicating that all the specimens exhibit a hardness peak above the quenched hardness value at approximately 525 °C.The hardening phenomenon during tempering is attributed to the decomposition of austenite, precipitation of supersaturated carbon, and precipitation of secondary phases.The hardness begins to diminish above 525 °C and steeply declines during the tempering stage between 600 °C and 650 °C, with a steeper decrease in hardness at higher hardness values.After reaching 650 °C, the hardness generally decreases to the same level.In general, martensitic structures fully decompose during tempering at approximately 450 °C, thus continuously decreasing hardness during tempering.However, the experimental steel begins to significantly soften above 600 °C, indicating the significant role of its alloy composition in the tempering stability of martensite.M50NiL, which is also used as a bearing steel, can exhibit a similar hardness pattern.After carburizing, the peak hardness is reached after quenching at 1050 °C and tempering at 525 °C.The highest hardness during carburization is 650 HV, and the core hardness is 410 HV.The hardening method is a secondary hardening mechanism caused by the precipitation of carbides during tempering [3].The other type, CSS-42 L, is a third-generation bearing steel with a significantly reduced hardness.After tempering at 500 °C, the hardness is 293.2HV.After tempering at 540 °C, the hardness is 366.7 HV.Finally, after tempering at 580 °C, the hardness is 453.8HV.Although it conforms to the law of secondary hardening, the hardness is greatly reduced [20].
DSC is employed to detect heat effects during the tempering process, helping determine the temperature and intensity of the second phase precipitation.During heating, the normal endothermic effect produces a significant heat capacity effect, obscuring the thermal effects of residual austenite decomposition and carbide  precipitation.The fully annealed specimens are utilized as reference samples, and the quenched specimens are used as test samples.The microstructures of the quenched specimens mainly consist of metastable martensite and residual austenite.Tempering results in decomposition and generates certain heat effects.The fully annealed specimens have equilibrium microstructures without martensite or residual austenite, and the precipitation and growth of carbides are relatively sufficient.Hence, no heat effects other than changes in heating capacity occur during tempering due to increases in temperature.
At a heating rate of 5 °C min −1 (figure 6(b)), an exothermic peak associated with residual austenite decomposition appears at 313 °C [21].The decomposed carbon first diffuses toward the matrix and then begins to precipitate carbides after 330 °C.The first carbide precipitation peak occurs at approximately 527 °C, corresponding to the peak hardness of the material.Subsequently, the carbides gradually grow and cause the hardness to decrease.At 590 °C, a clear endothermic peak appears, during which the alloy carbides continue to grow through long-range diffusion of the alloying elements.The coherence of the carbide interfacial plane is destroyed, and the hardness decreases catastrophically, resulting in the transformation of the microstructure to tempered sorbite.During the tempering process, an endothermic peak associated with austenite transformation appears at 742 °C, coinciding with Ac1 in the expansion curve.At 901 °C, an endothermic peak is observed due to the dissolution of alloy carbides and intermetallic compounds.
Compared with the heating curve at 5 °C min −1 , the heating rate of 20 °C min −1 is faster.In addition, there is no obvious peak temperature during residual austenite decomposition.The first peak appears at carbide precipitation at 584 °C, with a relatively sharp peak shape and obvious heat effect.This phenomenon arises due to the thermal lag caused by the fast heating rate.As the heating rate increases, the starting temperature is delayed, but the reaction becomes increasingly intense.Based on the peak shape, the exothermic peak at 753 °C corresponds to the carbide precipitation peak at 667 °C, and the endothermic peak at 835 °C corresponds to the austenite transformation peak.DSC is usually used for thermodynamic analysis of precipitates.JMAK equation is used for data analysis under isothermal conditions.Kissinger formula is used for data analysis under linear heating conditions [22].In this paper, linear heating is used, so the latter is used.For the non-isothermal transformation kinetic equation, the general formula is as follows [23]: ) and G(T(t)) are the nucleation and the growth rates, T the temperature and g a geometrical factor that depends on the dimensionality of growth.For an interface-controlled growth n is an integer; for a diffusioncontrolled growth n takes either integer or half-integer values.The formula obtained by simplifying Formula (1) is the commonly used Kissinger formula, which is based on E g β/T p , E g β/RT p , nE g /T p and 2nE g /RT p »2: Where T p is the temperature of the exothermic peak in the DSC curve, and R is the gas constant, β is the heating rate.E g is the activation energy for growth solely.Generally, the phase transformation and precipitation processes in materials meet these conditions.By substituting the peak temperature Tp of the exothermic peak into formula (2), the effective activation energy of the precipitate corresponding to the exothermic peak in the linear heating process of the material can be calculated.The experimental data are shown in table 2: From the slope E g /R of Kissinger curve obtained from ln(T p 2 /β) to 1/T p , it can be calculated that the effective activation energy during carbide precipitation at 527 °C was Ec1 = 53.02kJ mol −1 .During the transformation to tempered sorbite at 590 °C, the activation energy was Ec2 = 55.54 kJ mol −1 .At 667 °C, the activation energy was Ec3 = 112.85kJ mol −1 .During the austenite transformation at 742 °C, the activation energy was Ec4 = 121.68kJ mol −1 .

Type of carbide after tempering
After tempering the material, carbides are obtained through electrolytic extraction, and x-ray diffraction (XRD) is utilized to analyze the carbides.As shown in figure 7, after tempering at 500 °C, the primary carbides are VC and M 7 C 3 , and the number of M 23 C 6 carbides is not significant.Carbides are uniformly distributed in the microstructure as small spherical particles with diameters less than 0.1 μm between the matrix plate bundles (figure 8), and they are not concentrated or biased at the grain boundaries.After tempering at 550 °C, the M 7 C 3 diffraction peak does not significantly change, but the M 23 C 6 carbide diffraction peak appears, indicating that a significant number of M 23 C 6 carbides precipitates between 500 °C and 550 °C.The carbon atoms are redistributed among the alloying elements, and the atomic ratio changes.The matrix contains tempered martensite, and the ferrite begins to recover.After tempering at 550 °C, the VC carbide diffraction peak becomes less prominent, and the intensities of the M 23 C 6 and M 7 C 3 carbide diffraction peaks significantly increase.The matrix structure becomes more uniform, and the plate characteristics of the original martensite structure gradually weaken.The alloying elements diffuse over long distances and precipitate many irregularly shaped carbides of various sizes at the grain boundaries of the ferrites.Finally, after tempering at 600 °C, the VC diffraction peak almost disappears, and the M 23 C 6 carbide content increases.
During the tempering of Hybrid Steel 60, M 7 C 3 carbides preferentially precipitate due to the presence of chromium, which is a strong carbide-forming element with a high carbon content.According to figure 7, M 23 C 6 has not yet formed during tempering at 450 °C to 500 °C, and the increase in hardness occurs due to secondary hardening caused by the precipitation of M 7 C 3 carbides.As the carbides are much harder than the matrix, the uniform precipitation of spherical carbides with diameters less than 0.1 μm in the early stage can play a role in pinning grain boundaries and inhibiting the propagation of microcracks, thereby improving the fatigue performance of the material [24].After tempering at 500 °C, tempered martensite decomposition and ferrite recovery cause large accumulations of C and Cr at the grain boundaries, forming many carbide nucleation sites.Additionally, the newly formed carbides are primarily M 23 C 6 carbides with significantly lower carbon contents than the former carbides.Depending on the different matrix phases, carbon contents, and heat treatment processes, M 23 C 6 carbides can form inside austenite grains or at grain The carbides initially precipitate in the dislocation area and maintain a coherent relationship with the α phase.As the carbide content increases and the carbides grow and precipitate, the coherent strain between the carbides and the α phase becomes more pronounced until the hardness peaks [25][26][27].
As the tempering temperature increases, the carbide size increases, and the coherent relationship is disrupted.Next, the coherent strain disappears, and the dislocation density decreases.Full width at half maximum (FWHM) is directly associated with dislocation density, i.e. peak broadening is due to increase in dislocation density [28,29].At temperatures between 500 °C-550 °C, the FWHM variation is not pronounced.This observation suggests that carbides primarily precipitate in their nascent form, with a relatively minor accumulation of dislocations.Conversely, at temperatures spanning 550 °C-600 °C, the FWHM experiences a conspicuous increase.During this phase, carbides predominantly precipitate in an enlarged state, accompanied by a substantial accumulation of dislocations.
According to figure 8, we can observe the morphologies of the carbides, which are primarily spherical.Furthermore, from figure 8(b), it is evident that VC appears as dark-colored particles under the electron microscope, while carbides rich in chromium are primarily composed of bright-colored particles.According to the energy-dispersive x-ray spectroscopy (EDS) spectra for the three different states, the VC content is relatively low and remains consistent.However, the content of chromium-rich carbides gradually increases, which is in agreement with the changes in carbide content shown in figure 7.
According to figure 9, at 500 °C, the maximum diameter of the carbides is approximately 0.5 μm.Approximately 3% of the carbides have this diameter, while carbides with a diameter of 0.1 μm account for approximately 70%.At 600 °C, the carbide significantly increases, and its shape and distribution become uneven, with a maximum diameter exceeding 1 μm, accounting for approximately 4%.The proportion of carbides with a diameter of 0.1 μm decreases to approximately 50%.At this stage, the carbides can detrimentally affect the fatigue properties of the material.To eliminate this effect, an isothermal treatment can be performed to spheroidize the carbides.Spheroidized carbides can substantially avoid becoming the starting point of cracks, thereby improving the fatigue strength and fracture toughness of the material.V precipitates as carbides between 500 and 700 °C in bainite and martensite and dissolves at 800 °C [30][31][32].The VC diffraction peak in figure 7 To investigate the precipitation behavior of carbides, the relationship between carbide precipitation and interface binding strength was analyzed using transmission electron microscopy.As depicted in figure 10(a), after tempering at 500 °C, no M23C6 particles were discerned in the field of view, while the diameter of M 7 C 3 particles had exceeded 0.1 μm.Furthermore, figure 10(b) reveals conspicuous dislocations along the ferrite boundaries, suggesting, in accordance with the form of carbide precipitation, that the precipitation temperature  of M 7 C 3 particles should be below 500 °C, with a plane spacing of approximately 0.35 nm. Figure 10(c) indicates that at 550 °C, M23C6 particles have initiated precipitation, with a diameter less than 0.8 micrometers.Observations in figure 10(d) suggest that certain interfaces between carbides and the matrix exhibit no apparent dislocation lines, with relatively close plane spacings.M 23 C 6 particles measure 0.23 nm, and the ferrite spacing is 0.2 nm, indicating that carbides are in the early stages of precipitation, with the matrix in a semi-coherent state.Upon elevating the tempering temperature to 600 °C and observing the presence of M 23 C 6 particles, figure 10(e) reveals the formation of distinct dislocation lines between carbides and the matrix, with a significant increase compared with 550 °C, aligning with the FWHM results.At this juncture, the quantity of carbide particles within the field of view significantly increases, primarily distributed along grain boundaries.

Conclusions
(1) The expansion curve revealed Ac1 and Ac3 temperatures of 786 °C and 864 °C respectively for Hybrid Steel 60.The austenitization activation energy was calculated as 121.68 kJ mol −1 through the DSC curve.The tempered sorbite transformation temperature was 590 °C, exhibiting an activation energy of 55.54 kJ mol −1 .Within solid solution temperatures spanning from 900 °C to 1000 °C, an increase in temperature resulted in the growth of austenitic grains, a rise in residual austenite content, and a decline in hardness.Beyond 1050 °C, carbides commenced dissolution, subsequently elevating hardness.
(2) The effect of quenching temperature on material hardness was negligible, whereas the cooling rate during quenching had a significant impact on hardness.Hardness increased with the cooling rate during slow cooling, reaching a maximum at 5 °C s −1 , and then decreased with further increases in cooling rate.Consequently, after tempering, hardness proportionally escalated with post-quenching hardness.
(3) The residual austenite decomposition temperature during tempering approximates 313 °C.Following tempering, hardness escalated and subsequently declined with tempering temperature, attaining a zenith of 566 HV at approximately 525 °C.XRD analysis indicated that the extracted carbide primarily comprises M 7 C 3 and VC, exhibiting diameters below 0.1 μm in a spherical configuration at 500 °C.Beyond 500 °C, and M 23 C 6 started to precipitate and amass extensively in an irregular formation at the grain boundaries.The material's secondary hardening predominantly originated from the substantial precipitation of M 23 C 6 carbides.
(4) When carbides initially precipitate, they exhibit a coherent relationship with the matrix.However, as the carbides grow, this coherence is disrupted, leading to the continuous accumulation of dislocations.The interplanar spacing of the matrix is approximately 0.2 nm, while that of M 7 C 3 is around 0.35 nm, and M 23 C 6 is approximately 0.23 nm.

Figure 1 .
Figure 1.(a) Pretreatment of samples a, b, c, d; (b) Temperature rise of samples a, b, c, and d in a differential thermal analyzer.

Figure 2 .
Figure 2. (a) Expansion curve of mixed steel 60 when heated to 920 °C at 5 °C s −1 and cooled at 0.2 °C s −1 ; (b) Material hardness and residual austenite content at different quenching temperatures.

Figure 4 .
Figure 4. Optical micrographs after holding at 920 °C for 1 h and then water-cooled quenching.

Figure 5 .
Figure 5. Temperature of martensite transformation point at different cooling rates after holding at 920 °C for half an hour.

Figure 6 .
Figure 6.(a) Hardness curve after tempering at different cooling rates; (b) DSC curve of tempering process.

Figure 7 .
Figure 7. XRD spectra of carbide extracted at different temperatures.

Table 1 .
Chemical composition of test steel, wt%.