Evolution of Microstructure and the Influence of Carbides on Hardness Properties in Martensitic Stainless Steel 90Cr18MoV During Heat Treatment

Abstract

In this study, we utilized Thermo-Calc software (2023a) to optimize the heat treatment process of martensitic stainless steel 90Cr18MoV through phase diagram calculations. The microhardness of 90Cr18MoV was characterized using a nanoindentation instrument. The microstructural morphology of the samples was analyzed using scanning electron microscopy (SEM). The composition of the samples was characterized through scanning electron backscatter diffraction (EBSD) and X-ray diffraction (XRD). Additionally, laser confocal microscopy (FIB) and transmission electron microscopy (TEM) were employed to characterize the precipitate phase composition and size before and after heat treatment, while also observing the dislocation structure within the samples. The relationship between the quenching temperature and the percentage of residual austenite content in the material was established. The influence of the dislocation structure and precipitate size on the hardness of the samples was investigated. The research findings confirm that the observed secondary hardening phenomenon in tempered samples is attributed to the co-precipitation of two types of carbides, M₂₃C₆ and MC, within the matrix. The study investigated the effects of the tempering temperature and duration on the size of secondary precipitates, indicating that M₂₃C₆ and MC particles with sizes less than or equal to 20 nm contribute to enhancing the matrix, while particles larger than 30 nm lead to a reduction in hardness after tempering. Notably, during the tempering process, M₂₃C₆ precipitated from the matrix nucleates on MC.

1. Introduction

The high-carbon, high-chromium martensitic stainless steel 90Cr18 is widely used in the production of high-end bearings, surgical instruments, and valves due to its exceptional hardness and wear resistance [1]. The addition of alloying elements in high-carbon martensitic stainless steels can obstruct slip planes, thereby enhancing their mechanical properties [2]. Furthermore, an increase in the chromium content present significantly improves the material’s corrosion resistance [3,4]. Due to the strong reducibility of chromium (Cr), it is widely used in stainless steel formulations [5]. Similarly, the addition of a small proportion of molybdenum (Mo) has been shown to improve the resistance of stainless steel against pitting and crevice corrosion in chloride environments [6].

The alloy 90Cr18MoV is classified as a high-carbon, high-chromium martensitic stainless steel. Martensite forms as a supersaturated solid solution of carbon (C) within alpha iron (ɑ-Fe), making it easier for high-carbon steels to develop a robust martensitic structure [7]. During the melting and solidification stages of 90Cr18 martensitic stainless steel, there is a tendency for the formation of difficult-to-remove network or plate-like carbides such as M₇C₃ and M₂₃C₆; “M” mainly refers to Cr [8]. The formation of large-sized carbides like Cr₇C₃ and Cr₂₃C₆ inevitably leads to regions deficient in chromium surrounding these carbides, which significantly diminishes the localized corrosion resistance of 90Cr18 martensitic stainless steel. Moreover, atoms such as manganese (M) or vanadium (V) can substitute chromium atoms at positions within Cr₂₃C₆. This substitution releases Cr from the carbides into the matrix material, which greatly enhances localized corrosion resistance while also facilitating dislocation accumulation that contributes positively to increasing the material strength [9].

Early studies have indicated that the size and composition of M_xC_y carbides are related to the cooling rate, and an excessively large carbide size significantly affects material properties [10,11]. It is well known that heat treatment methods can facilitate the phase transformation from M₇C₃ to M₂₃C₆ [10]. The phase transformation may initiate either from the center of primary carbide M₇C₃ [12] or from the peripheral regions of M₇C₃ [10,13,14]. Wang et al. [15] discovered that in high-performance microalloyed (HP-MA) steels, the transition from M₇C₃ to M₂₃C₆ at a heat treatment temperature of 1000 °C (Quenching) is controlled by the diffusion of carbon from the shell of M₂₃C₆ into the matrix. Yu et al. [16] investigated carbide evolution in high-carbon 8Cr13MoV stainless steel and found it challenging to eliminate the formed network carbides during solidification under conventional heat treatment and thermal deformation techniques. However, thermal deformation proves to be effective in fragmenting network carbides to smaller sizes by controlling deformation temperatures [16]. The thermal deformation at lower temperatures than Ac1 yields better fragmentation results [17]. Nevertheless, the higher concentrations of C and Cr in 90Cr18 make it more feasible for the formation of coarse network carbides, leading to an increased hardness that could not to be worked cold. Consequently, the current processing of 90Cr18MoV martensitic stainless steel predominantly employs hot rolling techniques. Currently, the temperature of hot rolling processes is lower than 800 °C for energy saving, which could not completely dissolve and fragment these network carbides. Cracking inevitably occurs, requiring further toughening treatments. To investigate suitable heat treatment processes, Qian et al. [18] carried out spheroidizing annealing and quenching-tempering processes on 60Cr16MoMA martensitic stainless steels and found that the coarse, fragmented M₇C₃ carbides transformed into M₂₃C₆ carbides.

However, it was found that post-heat-treated 60Cr16MoMA did not eliminate the detrimental effects of coarse M₂₃C₆ on the material; instead, an increase in the size of M23C6 was observed after tempering, indicating a lack of precise control over carbide dimensions. The carbide is inevitably formed in high-carbon martensitic stainless steels. Large-sized carbides lead to decreased corrosion resistance and increased brittleness in materials. In comparison, fine carbides tend to enhance material properties such as the presence of nanoscale chromium-rich M₂₃C₆ carbides that enhance the hardness of the microstructure [19].

In recent years, there has been growing interest among researchers regarding the size and type of secondary precipitated carbides due to their beneficial role in strengthening the matrix [20]. Li et al. [21] indicated that large-sized M₂₃C₆ carbides with sizes ranging from 0.3 µm to 1.5 µm significantly improve low-temperature toughness in martensitic heat-resistant steels. Bonagani et al. [22] conducted a tempering process on the steels containing 13%wt. Cr at temperatures above 550 °C and found that the Cr-depleted zones with the sizes of approximately 7–9 nm appeared around the secondary precipitated carbides, resulting in a decline in the pitting corrosion resistance of the materials.

Currently, research on high-carbon martensitic stainless-steel carbides primarily focuses on the types of primary carbides formed in the martensitic matrix after heat treatment. However, there remains a lack of comprehensive studies regarding the types of secondary carbides and the size variations of these carbides at different temperatures. In this study, the influence of different quenching temperatures on the phase transformation of 90Cr18MoV was investigated, as well as the effect of the residual austenite content resulting from quenching on the hardness of the matrix and the whole materials. After quenching and tempering treatment, the mechanism of secondary hardening in 90Cr18MoV during the tempering process was investigated. By analyzing the size of secondary phase precipitates generated during tempering, we studied the influence of carbide dimensions on the hardness of the 90Cr18MoV microstructure. Additionally, we examined the relationship between the size of secondary precipitates and the tempering temperature. Furthermore, an analysis was conducted on the precipitation and formation mechanisms of M₂₃C₆.

2. Materials and Methods

2.1. Materials

The material utilized in this study is martensitic stainless-steel plate 90Cr18MoV, which has undergone hot rolling and annealing processes(This material has been independently smelted and rolled by our research institution.). The chemical composition of the steels is shown in

Table 1. Chemical composition of the 90Cr18MoV steel (=wt. %).

C	Si	Mn	P	S	Cr	Ni	Mo	V
0.903	0.512	0.420	0.026	0.004	17.19	0.329	1.069	0.161

. The sample size is 10 mm × 8 mm × 4 mm (length × width × height). The heat treatment process for 90Cr18MoV stainless steel is illustrated in Figure 1. As shown in Figure 1, this experiment investigates the quenching treatment of 90Cr18MoV within a temperature range of 1020 °C to 1100 °C, with increments of 20 °C as individual research subjects. The quenching duration is set at 20 min, and oil is utilized for cooling the samples. To explore the effects of different tempering temperatures on the internal structure and properties of the material, we examine a temperature range from 150 °C to 700 °C, stepping in increments of 50 °C as separate research subjects. The tempering time is established at 180 min. We designate the material after quenching as Q-W, where Q represents the quenching process and W denotes the quenching temperature. We designated the samples that underwent tempering as QT-X/Y, where QT denotes quenching followed by tempering, X represents the tempering temperature, and Y indicates the duration of tempering.

To investigate the microstructure of the obtained samples, the samples after heat treatment were polished and then etched with a mixture of 90 mL of 4% ferric chloride solution and 10 mL of hydrochloric acid for 5 s. After that, an electrolyte polishing process was performed on the samples at a voltage of 30 V and a current of 1 mA for electron backscatter diffraction analysis. The electrolyte solution contained 10% perchloric acid and 90% alcohol to obtain specimens suitable for EBSD analysis.

2.2. Characterization

The microstructure of 90Cr18MoV was observed using Leica-DMI8C optical microscopy (OM) (Wetzlar, Germany) and a TESCAN/CLARA GMH (Brno, Czech Republic) field emission scanning electron microscope (FESEM). An Oxford Symmetry S2 electron backscatter diffraction system (EBSD) and Oxford UltimMax65x energy-dispersive X-ray spectroscopy (EDS) (Abingdon, UK), in conjunction with a Thermo Scientific Talos F200X S/TEM (Waltham, MA, USA) transmission electron microscope (STEM), were used for electron microscopy observation. Transmission specimens were prepared utilizing a dual-beam thinning instrument and Helios 5UX focused ion beam (FIB) (Waltham, MA, USA) technology.

Nano-hardness and elastic modulus characterization of the samples were performed using a Bruker Nano Inc. (Billerica, MA, USA). TI Premier nanoindentation tester, applying a load of 10 mN.

3. Results

3.1. Quenching Treatment

Figure 2a shows the morphology of the microstructure of the 90Cr18MoV samples as received. The carbides (orange marker) occupied a large portion of the 90Cr18MoV matrix. The corresponding TEM images of carbides are shown in Figure 2b–d. The diffraction patterns indicate that these large carbides are composed mainly of M₇C₃ inside grains (marked in yellow) while the small carbides are composed mainly of M₂₃C₆ (marked in blue).

Martensite formed due to the post-quenching processes, as a supersaturated solid solution of carbon, has been believed to present high hardness [23]. To obtain high hardness, the 90Cr18MoV were quenched at different temperatures of 1020 °C, 1040 °C, 1060 °C, 1080 °C, and 1100 °C. We designated these samples as Q-1020, Q-1040, Q-1060, Q-1080, and Q-1100, respectively. The microstructure of 90Cr18MoV before and after different quenching temperatures are shown in Figure 3. The presence of a large area of precipitated carbides in the original sample, as shown in Figure 3a, is attributed to the annealing process that facilitates the precipitation within the carbide matrix, thereby softening the material and making it easier to process.

As the temperature increases, the carbides in the matrix gradually diminish. Larger carbides decrease in size, while smaller carbides completely dissolve back into the matrix. In Figure 3e, it can be observed that no minute carbides are present in the matrix of the Q-1080 sample. However, as shown in Figure 3f, with a further elevation of the temperature, carbides begin to reappear in the matrix of the Q-1100 sample. At the same time, an increase in carbon concentration of the matrix resulted in a decrease in the martensitic transformation temperature. This phenomenon leads to an incomplete austenite to martensite transformation under different quenching conditions. The phase structure analysis was conducted on the samples subjected to quenching using EBSD, and the results are presented in Figure 4. At a quenching temperature of 1040 °C, the retained austenite having an fcc structure is labeled in red and its area fraction is 0.5% as shown in Figure 4a–d. The fraction of the fcc structure increased to 1.15% after quenching at a temperature of 1060 °C. The fraction of the fcc structure increased to 6.42% after quenching at a temperature of 1080 °C. After quenching at 1100 °C, the fraction of the fcc structure reached up to 35.8%. Therefore, as the quenching temperature increases, the content of austenite in the matrix also rises. Concurrently, fine carbides gradually dissolve back into the matrix with increasing quenching temperatures, leading to an increase in carbon content within the matrix. Since carbon (C) is a stabilizing element for austenite [24], a higher carbon content enhances the stability of austenite during heat treatment. Consequently, residual austenite becomes more prevalent with rising quenching temperatures, which aligns with findings from other studies [25,26]. In this research, we determined that the critical quenching temperature for carbon dissolution is 1060 °C.

To investigate the effect of an increasing quenching temperature on the microstructure and hardness of 90Cr18MoV, we conducted hardness characterization on samples quenched at different temperatures, as illustrated in Figure 4d,e. Under identical quenching times, the macroscopic hardness of 90Cr18MoV initially increases with a rising quenching temperature before subsequently decreasing. A similar trend is observed in the microscopic organizational hardness, which also shows an initial increase followed by a decline. Notably, samples quenched at 1080 °C exhibit maximum hardness values of 61 HRC and 8.8 GPa. The hardness of 440C martensitic stainless steel obtained through deep cryogenic treatment, as reported by A. Idayan et al., is approximately 60 HRC [27]. To validate that the measured microscopic hardness reflects the material’s structural integrity, we performed SEM scanning analysis on samples post-nanoindentation testing, as shown in Figure 4f.

3.2. Tempering Treatment

To eliminate the internal structural defects of samples after rolling and quenching, it is necessary to perform tempering treatment on the quenched samples. As shown in Figure 4, the hardness of samples quenched at a temperature of 1060 °C is the highest. Additionally, from the EBSD results of the quenched samples, it can be observed that there is minimal residual austenite present when quenched at this temperature. Therefore, we will perform wire cutting on the Q-1060 samples to divide them into multiple smaller specimens for subsequent tempering treatment.

As illustrated in Figure 5, SEM scanning images of tempered samples at different temperatures reveal that the microstructure primarily consists of martensite and some dispersed carbides. Furthermore, with an increase in the tempering temperature, a significant number of precipitates appears within the matrix structure. Specifically, as depicted in Figure 5a, numerous precipitates are observed in the matrices of samples tempered at temperatures of 550 °C and 600 °C, respectively. Zhenfeng Xu et al. [28] conducted a study on the precipitation behavior of M₂₃C₆ under various boundary conditions, confirming that M₂₃C₆ precipitates at a temperature of 600 °C.

To investigate the relationship between the hardness of tempered samples and tempering temperature, we conducted a macro-hardness analysis on samples tempered at different temperatures, as illustrated in Figure 5b. The hardness of the samples exhibited an initial increase followed by a decrease with rising tempering temperatures. Additionally, to examine the secondary hardening characteristics of the matrix in tempered samples, we performed nanoindentation hardness analysis on specimens tempered at 460 °C, 480 °C, 500 °C, and 520 °C. By combining nanoindentation techniques with scanning electron microscopy (SEM), we assessed the hardness of the matrix for these specific tempering temperatures. As shown in Figure 5c,d the maximum nanoindentation hardness achieved during secondary hardening of the tempered samples was found to be 7.2 GPa. The maximum macrohardness is 57 HRC.

4. Discussion

4.1. Phase Diagram Calculation

The phase diagram was calculated using Thermo-Calc software, as illustrated in Figure 6a. The high-temperature liquid phase exists above 1250 °C, while the austenite phase is present within the temperature range of 800 °C to 1380 °C. In this study, the quenching temperatures employed are 1020 °C, 1040 °C, 1060 °C, 1080 °C, and 1100 °C; Thus, these quenching temperatures fall within the austenite phase region. Furthermore, with an increase in the quenching temperature, the relative content of the austenite phase also gradually increases, which is consistent with the EBSD results presented in Figure 4. The bcc structure appears between temperatures of 1350 °C and 1420 °C, as well as below temperatures of 830 °C. The two-phase region comprising fcc and bcc structures exists within the temperature range of 800 °C to 850 °C. Therefore, all tempering processes conducted in this study occur outside this two-phase region. M₂₃C₆ can be found at temperatures below 1180 °C while M₇C₃ may appear within a temperature range from 1150 °C to 1270 °C.

Figure 6b presents the phase composition diagram of M₂₃C₆ at equilibrium, simulated using Thermo-Calc. It is evident that there is a significant increase in the presence of M₂₃C₆ within the temperature range of 330 °C to 385 °C, after which it begins to re-dissolve into the matrix at temperatures exceeding 800 °C. Although the processing methods employed in this study do not allow for the samples to achieve complete equilibrium at identical temperatures, the overall trend remains consistent. Consequently, during the tempering process, precipitation of M₂₃C₆ will occur. In order to investigate the precipitation conditions of the M₂₃C₆ phase that are compatible with industrial production, we will employ transmission electron microscopy for a more detailed analysis of the samples.

To investigate the precipitation models of M₂₃C₆ and MC at different temperatures, this study employs Thermo-Calc to analyze and simulate the nucleation and growth of M₂₃C₆ and MC in samples using both bulk nucleation and grain boundary nucleation models. As shown in Figure 6c, under isothermal conditions at 500 °C, simulations were conducted to assess the nucleation and growth of M₂₃C₆ over varying holding times in a fully equilibrated state. Similarly, Figure 6d presents simulations for the same parameters at 600 °C.

It can be observed that at a temperature of 500 °C, as the holding time increases, both M₂₃C₆ exhibit continuous growth in their nucleation sizes. Likewise, at a temperature of 600 °C, an increase in the holding time also leads to an enhancement in the size of M₂₃C₆ nuclei. However, it is noteworthy that the sizes of precipitated carbides (M₂₃C₆) formed at 600 °C are significantly greater than those formed at 500 °C.

4.2. Microstructural Characterization

In order to investigate the reasons for the changes in hardness of samples under different tempering conditions, we conducted STEM analysis on samples tempered at various temperatures, as illustrated in Figure 7. Figure 7a–c represent the tempered samples at temperatures of 150 °C, 250 °C, and 350 °C, respectively. It can be seen that the number of dislocations formed after quenching significantly decreased with an increase in the tempering temperature. Zhang et al. [29] investigated the toughening mechanisms during the tempering process of high chromium steel, revealing that as the tempering temperature increases, the dislocation structure within the experimental steel significantly diminishes. Combining this observation with the relationship curve between the sample hardness and tempering temperature shown in Figure 5, we conclude that higher tempering temperatures effectively reduce structural defects; consequently, as these organizational defects diminish, the material hardness also tends to decrease. It is evident that as the tempering temperature increases, the number of residual structural defects (dislocations) remaining after quenching significantly decreases. This microstructural morphology corroborates the relationship curve between sample hardness and tempering temperature presented in Figure 5.

However, a continuous increase in the tempering temperature resulted in the formation of nanoscale precipitates within the matrix of the 90Cr18MoV alloy as shown in Figure 8a–c. Wang et al. [30] conducted a study on the aging strengthening of high-carbon high-Cr steel and found that the carbides precipitated during the tempering process have a certain reinforcing effect on the matrix. Furthermore, accumulation of dislocations around these fine carbides was observed, leading to a secondary hardening effect observed in the tempered samples. To gain a deeper understanding of the specific characteristics of these precipitates, further detailed analyses are required.

To further investigate the lattice structures of the above samples, we initially employed Electron Backscatter Diffraction (EBSD) to study the lattice structure of the samples post-tempering; the results are shown in Figure 9. Figure 9a presents the QT-500/180. The results indicate the presence of smaller carbides surrounding larger carbides. These are small carbides, specifically MoC and VC; some researchers have experimentally verified that carbides such as VC precipitate during the tempering process [31,32]. They are observed to precipitate exclusively around the particles of the larger carbides. The tempered samples were scanned by XRD and are shown in Figure 9c. The results again verified the existence of MoC and VC in the matrix after quenching. Furthermore, it was noted that the content of M₂₃C₆ and MC increased continuously with rising tempering temperatures.

To verify the specific structure and composition of carbides precipitated during high-temperature tempering, TEM diffraction analysis was conducted on the carbide precipitation regions. FIB was employed to cut and thin the tempered samples as shown in Figure 10a,b. Figure 10c presents the QT-450/180 in which no fine carbides were observed within the matrix. In contrast, Figure 10d illustrates QT-500/180 where the average size of the precipitated carbides is approximately 20 nm. The corresponding diffraction pattern is shown in Figure 10e. It indicates that the carbides are VC and M₂₃C₆. The positional correlation of the diffraction spots indicates that M₂₃C₆ is nucleated on MC. Figure 10f illustrates the transmission dark field image of the precipitates in the QT-600/180 sample. The corresponding diffraction pattern is shown in Figure 10g, and the carbides were verified to be M₂₃C₆ and the average size of the precipitated carbides is 50 nm as shown in Figure 10f. Figure 10j shows QT-700/180 with the transmission dark field image. The corresponding diffraction pattern is shown in Figure 10k, and the carbides were verified to be M₂₃C₆ while the average size larger than 100 nm is shown in Figure 10j. These small M₂₃C₆ directly precipitate from the matrix during post-tempering and grow gradually with the increasing tempering temperature.

4.3. Refinement of Precipitate Size Statistics

The carbides precipitated during tempering are randomly dispersed within the matrix, contributing to dispersion strengthening. From Figure 10, it can be observed that the precipitated carbides exhibit an fcc structure corresponding to M23C6. To quantitatively investigate the effect of carbide size on the hardness performance of 90Cr18MoV due to tempering, we employed transmission electron microscopy (TEM) to statistically analyze the sizes of precipitates in samples tempered at different temperatures, as illustrated in Figure 10j, Figure 11b and Figure 12d,f. As the tempering temperature increases, the average size of the carbides gradually rises. Specifically, after tempering for 180 min at a temperature of 500 °C, the average size of precipitates is measured at 19.9 nm. At a temperature of 600 °C, the average size of carbides in tempered samples reaches 54.7 nm; while at a temperature of 700 °C, this value further increases to an average size of 152.86 nm. The overall trend is consistent with the simulation calculation illustrated in Figure 6.

At the same time, to investigate the effect of different tempering durations on the size of secondary precipitated carbides at a constant temperature, we conducted tempering experiments on samples that were quenched at 1080 °C for various durations (30 min, 60 min, 120 min, and 180 min) at a temperature of 600 °C. The results are illustrated in Figure 12a–e. The average size of the precipitates after tempering for 30 min at 600 °C was approximately 30 nm. After tempering for 60 min, the average size increased to about 38 nm; following a duration of 120 min, it reached around 41 nm; and after tempering for an extended period of 180 min, the average size further increased to approximately 54 nm. As shown in Figure 12e it is evident that with an increasing tempering time, there is a corresponding decrease in the sample hardness.

In the process of studying the effect of different tempering times at the same temperature on the size of precipitated carbides in samples, we observed that at a tempering temperature of 500 °C, the sizes of precipitates varied with different tempering durations, as illustrated in Figure 11a,b. The size of precipitates after 60 min of tempering was approximately 10 nm, while after 180 min it increased to around 20 nm. However, despite the increase in tempering time, there was no significant change in sample hardness within the margin of error; as shown in Figure 11c, all samples maintained a hardness level around 59 HRC. By correlating Figure 10j, Figure 11b and Figure 12d, we conclude that when the tempering temperature exceeds 500 °C, an increase in this temperature leads to continuous growth of secondary precipitate M₂₃C₆ within the sample matrix. When M₂₃C₆ is sized at or below approximately 20 nm, it contributes positively to enhancing the sample hardness. Conversely, once M₂₃C₆ exceeds a size greater than 30 nm, its further growth results in a declining trend in the sample hardness.

4.4. The Formation Mechanism of M₂₃C₆

The presence of M₂₃C₆ in the tempered samples leads us to conclude that, in the samples tempered at 500 °C, we observed the formation of carbides with sizes of 20 nm and below. The TEM images of the QT-450/180 samples revealed that such fine carbides were not present in the matrix (see Figure 10d), Therefore, it is certain that the fine carbides precipitated at 500 °C are formed because of the diffusion of C and Cr elements within the matrix. Additionally, the presence of VC was identified within these small carbides as shown in Figure 9. The XRD and EBSD analyses conducted on different tempering samples indicate that MC is present in the tempered samples, with its quantity increasing alongside the rise in the tempering temperature. Furthermore, the analysis of diffraction spots presented in Figure 10e demonstrates that the secondary precipitation of M₂₃C₆ during tempering occurs at MC as a nucleation point. The schematic diagram illustrating the nucleation of M₂₃C₆ on MC is shown in Figure 13. Furthermore, this precipitate phase is beneficial for facilitating the nucleation of the fine carbides within the matrix [33]. The combination of V with C occupies the sites that were originally held by Cr and C pairs. The released Cr elements deceased the amount of precipitated chromium-containing carbides and affected the corrosion resistance of samples. Moreover, the growth behavior of these carbide precipitates is influenced by elemental diffusion.

5. Conclusions

This study investigated the effects of quenching and tempering temperatures and durations on the microstructure and mechanical properties of 90Cr18MoV, leading to the following conclusions.

(1) The 90Cr18MoV alloy, quenched at a temperature of 1080 °C, presented a small amount of retained austenites. The increase in the carbon content within the matrix contributes to the increase in hardness. The measured microhardness and nano-hardness were 61 HRC and 8.8 GPa, respectively.

(2) The effect of the tempering temperature on the quenched 90Cr18MoV samples at a temperature of 1080 °C was investigated. A maximum microhardness and nanohardness of 57 HRC and 7.2 GPa were achieved at the tempering temperature of 500 °C due to the secondary hardening.

(3) The phenomenon of secondary hardening during tempering is attributed to the combined effects of the precipitation of MC and M₂₃C₆ carbides, which are smaller than 20 nm in size, within the matrix. The M₂₃C₆ that precipitates during tempering serves as a nucleation site for further nucleation processes occurring in the matrix, originating from MC.

(4) At a tempering temperature of 500 °C, maintaining the process for 180 min can result in the formation of secondary precipitates approximately 20 nm in size, which enhances the hardness of the microstructure. In contrast, secondary precipitates larger than 30 nm tend to reduce the hardness of the microstructural features.