Precipitation and Age-Hardening in Fe-25Co-15Mo Carbon-Free High-Speed Steel via Hot Isostatic Pressing

Abstract

High resistance to tempering and extended service life are pivotal research directions for cutting tools utilized in the machining of industrial machine tool. The design of alloys and their manufacturing processes have become methods for the development of cutting tool materials. Carbon-free Fe-Co-Mo steel (FCM) has garnered attention due to its excellent magnetic properties and high-temperature performance, as well as its superior thermal conductivity, making it an ideal choice for applications in high-temperature and high-pressure environments. The µ-phase within this alloy exhibits exceptional high-temperature stability and resistance to aggregation. Its characteristics suggest that it has the potential to replace carbide reinforcement phases, which are prone to coarsening, in high-temperature applications of powder high-speed steel. This application of the µ-phase could lead to an enhancement in the resistance to tempering and the service life of powder metallurgy high-speed steel cutting tools. However, there is a relative scarcity of published research regarding the preparation of carbon-free high-speed steel via hot isostatic pressing (HIP) technology and the subsequent heat treatment processes. In this study, Fe-Co-Mo alloys reinforced with the intermetallic compound µ-phase were prepared at hot isostatic pressing sintering temperatures of 1200 °C, 1250 °C, and 1350 °C. Furthermore, to investigate the influence of the solid-solution treatment temperature on the microstructure and macroscopic properties of the alloy, the as-prepared materials were subjected to solution annealing treatment at different temperatures (1120 °C, 1150 °C, 1180 °C, and 1210 °C). The results demonstrate that by moderately reducing the sintering temperature, the segregation phenomenon of the reinforcing µ-phase was significantly reduced, leading to an optimization of the microstructural uniformity of the prepared sample, with the micro-scale µ-phase being uniformly dispersed within the α-Fe matrix. As the temperature of the solid-solution annealing increased, the microstructural uniformity was further enhanced, accompanied by a reduction in the quantity of the reinforcing phase and refinement of the grain size. Notably, after solid-solution annealing at 1180 °C, the hardness of the samples reached a peak value of 500.4 HV, attributed to the decrease in the reinforcing phase and grain refinement during the annealing process. Aging treatment at 600 °C for 3 h facilitated the uniform precipitation of the nano-scale µ-phase, resulting in a significant increase in sample hardness to approximately 900 HV. The prepared material exhibited excellent resistance to tempering, indicating its potential for application in high-temperature service environments.

1. Introduction

Ordinary high-speed tool steel is renowned for its toughness and machinability, surpassing hard metals and ceramics. However, it has limitations in tempering and wear resistance. Despite being machined in an annealed state and hardened through high-temperature quenching and tempering, the inevitable martensitic phase transformation can cause distortion, increasing machining costs and complexity and potentially affecting the geometric precision of the machined parts. Thus, while offering benefits in workability and toughness, its application in high-temperature and wear-resistant scenarios is hindered by these issues [1,2,3,4].

With the growing demand for high-performance cutting tools, especially in the fields of aerospace and automotive manufacturing, greater requirements are placed on the material’s resistance to tempering, wear resistance, toughness, and machinability. Traditional high-speed steels contain high carbon content, which provides good hardness and wear resistance, but they are prone to softening at high temperatures, limiting their application under more severe conditions [5,6,7].

Carbon-free Fe-Co-Mo steel, as a new type of alloy, has attracted attention due to its excellent magnetic properties and high-temperature performance, and the superior thermal conductivity and toughness of FCM steel make it an ideal choice for cutting titanium alloys. Since the 1930s, many researchers have studied the preparation process, composition, heat treatment process, and strengthening mechanisms of FCM steel [8,9,10,11].

Compared to the traditional casting method, which has the problem of coarse carbide particles and uneven distribution, powder metallurgy can better control the microstructure of FCM materials, reducing the coarseness and segregation of carbides and thereby improving the performance of the material. With the development of hot isostatic pressing (HIP) technology, it can be used in the preparation of powder metallurgy Fe-Co-Mo carbon-free high-speed steel. High pressures (133–310 MPa) and high temperatures (up to 2500 °C) are applied simultaneously with inert gases (such as argon), densifying under high-temperature and high-pressure conditions and eliminating the pores and defects that may be produced during the powder metallurgy process, thereby further improving the material’s density and performance and improving the microstructure to obtain finer and more evenly distributed grains, which has a positive effect on improving the mechanical properties and toughness of FCM steel [12,13,14,15].

After solid solution treatment (>1000 °C) and rapid cooling, the (Fe,Co)₇Mo₆ intermetallic compound (μ-phase) will partially dissolve, and the hardness of FCM steel will relatively decrease. However, after precipitation hardening treatment at temperatures between 580 °C and 650 °C, the precipitation of nanoscale (Fe,Co)₇Mo₆ intermetallic compounds (μ-phase) will significantly increase the hardness. The Fe-25Co-15Mo alloy can reach a peak hardness of about 65 HRC and exhibit excellent temper stability [8,16,17,18].

Currently, research on Fe-Co-Mo carbon-free high-speed steel materials mainly focuses on the exploration and optimization of preparation processes and the improvement of the material’s comprehensive performance. The use of powder metallurgy technology allows for precise control of the material’s microstructure and chemical composition, resulting in a more uniform organizational structure and higher performance [18,19,20].

In this study, by using hot isostatic pressing technology at different temperatures, the aim is to suppress the coarsening and uneven distribution of grains in the FCM sintering process. By adopting solid solution annealing processes at different temperature, the appropriate temperature heat treatment process is obtained to further improve its hardness and wear resistance. This study systematically investigates the impact of different heat treatment processes on the material’s microstructure, nanoscale precipitation, and mechanical properties, exploring the preparation methods and performance of Fe-Co-Mo carbon-free powder high-speed steel and aiming to provide a material that maintains excellent performance at high temperatures [21,22].

2. Materials and Methods

2.1. Materials

Following the composition of Fe-25%Co-15%Mo (mass fraction), Fe-Co-Mo alloy powder was prepared using nitrogen gas atomization, with the raw material element purity reaching above 99.90%, and the powder particle size was controlled to be less than 175 μm. The mixed powder was loaded into a low carbon steel container and subjected to hot isostatic pressing treatment at temperatures of 1200 °C, 1250 °C, and 1350 °C. The pressure was set at 150 MPa, accompanied by a holding time of 3.5 h, in order to accomplish powder densification, as depicted in Figure 1. The heat treatment process of the sample was carried out as shown in Figure 2, after HIP treatment, samples with machined dimensions of 10 × 10 × 10 mm³ were taken for subsequent heat treatment process experiments.

The samples were annealed at 900 °C for 9 h in a vacuum heat treatment furnace (KJ-T1400-L1030WQ, ZKDL, Zhengzhou, China) to eliminate the internal stress and optimize the microstructure. In order to further improve the microstructure and the properties of the material, the sample was annealed in solution at 1120 °C, 1150 °C, 1180 °C, and 1210 °C and then put into the coolant for rapid cooling. Finally, the samples were precipitated and strengthened at 600 °C for 3 h.

2.2. Microstructure Characterization

In the current study, all test samples were obtained through electrical discharge machining (DK7735, NZMC, Ningbo, China) through wire cutting of the Fe-Co-Mo alloy in hot isostatic-pressed, solution-treated, and precipitation-hardened states.

In terms of performance testing and analysis, we first measured the relative density of the samples using a density measuring instrument according to the Archimedes principle to assess the degree of material densification, and the average measured density value was achieved by averaging the results from five individual tests to obtain a reliable measured density value. The relative density is the percentage ratio of measured density to theoretical density, used to quantify the compactness of as-HIP billet. Fe-25Co-15Mo material has a theoretical density of 8.35 g/cm³, and its calculation method is shown in formula one. Then, the microstructure of the samples was observed using an optical microscope (OM, Axioscope 5, Zeiss, Oberkochen, Germany) and a scanning electron microscope (SEM, Phenom XL G2, Phenom, Eindhoven, The Netherlands) to evaluate the size, shape, and distribution of the grains; compositional analysis (SEM with EDS detector, SEM5000X, CIQTEK, Hefei, China) of the samples was conducted to detect the segregation or precipitation of Fe, Co, and Mo elements.

$$Relative Density={\displaystyle \frac{{\rho }_{measured}}{{\rho }_{theoretical}}}\times 100\%$$

(1)

where ${\rho }_{measured}$ is the measured density and ${\rho }_{theoretical}$ is the theoretical density.

To qualitatively analyze the phase composition of Fe-Co-Mo alloy at different temperatures, the phase diagram was calculated using the JMatPro7.0.0 software, with an initial temperature of 600 °C and a simulation step size of 10 °C.

An X-ray diffractometer (XRD, D8 Advance, Bruker, Karlsruhe, Germany) with a Co target was used for phase characterization, with a test angle of 5–120° and a scanning speed of 5 °/min, to determine the existing phases and their distribution, thereby assessing the crystal structure of the material. A transmission electron microscope (TEM, JEM-2100, JEOL, Tokyo, Japan) was used to study the nano-scale strengthening phases. TEM foils were prepared by mechanically polishing the samples to a thickness of about 50 μm, followed by ion thinning. The microstructure at the nano-scale was observed.

2.3. Mechanical Tests

The Vickers hardness measurements (XHVT-1000Z, SCTMC, Shanghai, China) were performed on the polished surface of specimens with dimensions of 10 × 10 × 10 mm³, and 7 measurements were conducted to acquire the average data after excluding the maximum and minimum hardness data. The specimens of Fe-Co-Mo alloy in a hot isostatic state, a solution-treated state, a precipitation-hardened state, and a resistance to tempering test state were tested using the above method. The resistance to tempering test method involves subjecting the alloy in the precipitation-hardening state to a controlled argon atmosphere within a tube furnace at 600 °C. Subsequently, the alloy is removed and air-cooled to room temperature after specific holding times of 1 h, 3 h, 5 h, 10 h, 20 h, 40 h, 80 h, and 120 h.

3. Results and Discussion

3.1. Effect of Sintering Temperature on Microstructure and Mechanical Properties

Table 1. HIP temperature and material properties.

HIP Temperature (°C)	Density (g/cm3)	Relative Density (%)	Hardness (HV0.2)
1200	8.30	99.40	473.1 ± 5.6
1250	8.32	99.64	506.6 ± 6.2
1350	8.34	99.99	592.0 ± 6.5

shows the density and hardness of Fe-Co-Mo carbon-free high-speed steel under different hot isostatic pressing temperatures, and the density obtained by three different hot isostatic pressing temperatures is greater than 99%, indicating that the hot isostatic pressing process is feasible for the preparation of Fe-Co-Mo carbon-free high-speed steel. Within a certain temperature range, with the increase in the hot isostatic pressing temperature, the density of the material will be improved. The hardness of materials obtained at different hot isostatic pressing sintering temperatures is also different, and the hardness of materials sintered at 1350 °C is slightly greater than that after sintering at 1200 °C and 1250 °C.

In Fe-Mo-Co alloys, Co and Fe occupy the same sublattice, forming the µ-phase with Mo and enhancing the hardness of the alloy. Figure 3 shows the computational phase diagram of the Fe-Co-Mo alloy. Based on the phase diagram calculated using the JMatPro software for the Fe-Co-Mo alloy, we can qualitatively analyze the phase composition at different temperatures. The diagram illustrates the phase composition changes of the Fe-25Co-15Mo alloy at various temperatures. It can be observed from the diagram that as the temperature increases, the phase composition of the alloy undergoes significant changes. At 600 °C, the alloy is primarily composed of α-Fe and the μ-phase, with a higher content of α-Fe. As the temperature rises, the content of the μ-phase gradually increases while the content of α-Fe decreases. When the temperature reaches around 800 °C, the content of the μ-phase reaches its maximum value, and the content of α-Fe drops to its lowest. This indicates that within this temperature range, the μ-phase has higher stability, while the stability of α-Fe is lower. As the temperature continues to increase to 1000 °C, the γ-Fe phase begins to appear and gradually increases with the rise in temperature. This suggests that at higher temperatures, the stability of the γ-Fe phase increases while the stability of the μ-phase decreases. Eventually, when approaching 1400 °C, the γ-Fe phase becomes the dominant phase, and the contents of the μ-phase and the α-Fe phase almost disappear. This indicates that at this temperature, the stability of the γ-Fe phase is the highest. Li et al. discussed the Fe-Co-Mo ternary phase diagram, also involving alloys with compositions close to the one in the current study. They determined that the alloy contains a BCC phase and a μ-phase, and they established that the μ-phase acts as the strengthening phase in Fe-Co-Mo material [23].

Figure 4 shows the microstructure morphology of sintering at 1200 °C, 1250 °C, and 1350 °C for 3.5 h after hot isostatic pressing. According to Figure 4d,e, after sintering at 1200 °C and 1250 °C, it is evident that the white strengthening phase is uniformly dispersed within the gray-black matrix in the form of irregular particles and massive aggregates, respectively, the black spots observed in Figure 4 are likely areas where the strengthening phase, a hard and brittle intermetallic compound, has been selectively removed during the polishing process. After sintering at 1200 °C and 1250 °C, the particle size of the white particle strengthening phase is 2.94 μm and 2.92 μm, respectively, with minimal disparity between them. Based on Figure 4c,f, it can be observed that the distribution of the strengthening phase in the material after sintering at 1350 °C is irregular within the matrix, accompanied by a smaller particle size (1.95 μm) of the white strengthening phase as well, the high temperature can accelerate the dissolution of strengthening phases, which might then re-precipitate as the material cools down. This re-precipitation may not occur uniformly, resulting in a heterogeneous distribution of the light phase. The presence of grain boundaries and other microstructural features can act as preferential sites for the nucleation of new phases, contributing to the observed heterogeneity. Considering both the microstructure analysis and the uniformity of the strengthening phase distribution, a more suitable sintering temperature for Fe-Co-Mo alloy would be 1200 °C. The microstructure and properties of samples sintered at this temperature will now be analyzed.

The EDS analysis of the hot isostatic pressing treatment reveals the elemental distribution within the Fe-Co-Mo alloy, as depicted in Figure 5. This figure details the compositional analysis of two distinct regions, namely Spot1 and Spot2, along with both line and surface scanning analyses of the elemental distribution. The EDS data indicate a non-uniform distribution of elements across the alloy. In Spot1, the atomic percentages of Fe, Co, and Mo are measured as 42.82%, 24.78%, and 32.40%, respectively. In contrast, Spot2 shows a higher concentration of Fe at 61.26%, with comparatively lower percentages of Co and Mo at 33.98% and 4.76%, respectively. Figure 5g presents a line scan that demonstrates the fluctuation in the concentration of Fe and Mo, with Co maintaining a relatively constant level and exhibiting minor variations. Compared with Figure 5e,f, there is a difference of nearly 20% Fe content between the matrix phase and the strengthening phase, and the matrix phase has a higher Fe content. The difference between the bright and dark areas is mainly due to the large difference in signal contrast of the EDS probe. A higher count per second (CPS) value of Mo could signify the presence of the intermetallic µ-phase, whereas an increased CPS value of Fe might correlate with the formation of the α-Fe matrix phase. This uneven distribution of elements is presumably a result of the coexistence of two phases within the alloy.

Figure 6 shows the microstructure of Fe-Co-Mo alloy after hot isostatic pressing at 1200 °C × 3.5 h and annealing at 900 °C × 9 h. Figure 6a,b show the microstructure of the alloy at different magnifications, respectively. The strengthening phase is evenly distributed in the matrix, and it is found that the number of strengthening phases is significantly reduced upon comparison with Figure 4d. Through particle size statistics, the particle size of the strengthening phase is about 1.80 μm, which is significantly reduced compared with the particle size of the hot isostatic pressure state. During the annealing process at 900 °C, part of the strengthening phase may dissolve and enter the solid solution, resulting in a decrease in the amount of the strengthening phase. At the same time, high-temperature annealing may promote the growth of grain, but because of the dissolution of the strengthening phase, the grain growth is inhibited, resulting in the reduction of the particle size.

3.2. Effect of Solution Annealing Temperature on Microstructure and Mechanical Properties

Figure 7 shows the microstructure of Fe-Co-Mo alloys after sintering at 1200 °C × 3.5 h through hot isostatic pressing and annealing at 900 °C × 9 h at 1120 °C, 1150 °C, 1180 °C and 1210 °C, respectively. The white reinforcing particles are distributed evenly in the matrix. Solid solution annealing at different temperatures has a significant effect on the microstructure of Fe-Co-Mo alloy, especially the particle size and the quantity of the strengthening phase. Based on particle size statistics, with the increase in the solid solution annealing temperature, after solution annealing at 1120 °C, 1150 °C, 1180 °C, and 1210 °C, the particle size of the strengthening phase is 2.23 μm, 2.11 μm, 1.87 μm, and 1.66 μm, respectively. The particle size and quantity of the strengthening phase particles decrease gradually. In the process of solid solution annealing at a higher temperature, the strengthening phase may partially dissolve into the matrix and then undergo rapid cooling. High temperatures may promote the growth of grains, but due to the dissolution of the strengthening phase, the grain growth is inhibited, resulting in the refinement of the particles of the strengthening phase. In addition, the diffusion rate of elements increases at high temperatures, which may lead to the decomposition of the strengthening phase and the redistribution of elements. Thus, the particle size and quantity of the strengthening phase are affected.

Figure 8 shows the age-strengthening microstructure of Fe-Co-Mo alloy after treatment at different solid solution annealing temperatures (1120 °C, 1150 °C, 1180 °C, 1210 °C), and the white strengthening phase is uniformly distributed in the matrix in a granular form. With the increase in the solution annealing temperature, the particle size and quantity of the strengthening phase change. After solution annealing at 1120 °C, 1150 °C, 1180 °C, and 1210 °C, the particle sizes of the strengthening phases were 2.14 μm, 2.26 μm, 1.85 μm, and 2.00 μm, respectively. When the temperature rises from 1150 °C to 1180 °C, the particle size of the material decreases to a certain extent, and the particle size of the strengthening phase changes little compared with that of the solid solution annealed state in Figure 7, but the number of strengthening phases increase significantly. In the process of high-temperature solution annealing, the strengthening phase may partially dissolve into the matrix, and then the aging process will promote the redistribution or precipitation of solute atoms in the solid solution to form secondary phase particles.

The interaction between grain growth promotion and enhanced phase dissolution at high temperatures in Figure 7 and Figure 8 is complex. Although high temperatures may promote grain growth, dissolution of the strengthening phase inhibits this process, resulting in impeded grain growth. This dissolution behavior is reflected in the phase in Figure 3, which reveals changes in the phase stability of the alloy at different temperatures. At the same time, the increase in the element diffusivity at high temperature may lead to the decomposition of the strengthening phase and the redistribution of the elements, which will affect the stability and grain growth of the strengthening phase.

Figure 9 shows the grain size at different heat treatment stages of Fe-Co-Mo alloys. The experimental results show that the grain size of the alloy obviously increases with the increase in the heat treatment temperature. Specifically, after hot isostatic pressing at 1200 °C, 1250 °C, and 1350 °C, the grain sizes of the alloys are 1.26 μm, 1.45 μm, and 1.66 μm, respectively. Furthermore, the grain size of the sample treated at 200 °C was increased to 2.82 μm after annealing at 900 °C. After solution annealing at 1120 °C, 1150 °C, 1180 °C, and 1210 °C, the grain sizes of the samples increased to 4.81 μm, 4.93 μm, 6.01 μm, and 6.87 μm, respectively. After solution annealing, the samples were strengthened by aging at 600 °C, and the grain sizes were reduced to 2.39 μm, 2.52 μm, 2.77 μm, and 1.15 μm, respectively. The change in the grain size is closely related to the temperature and time of the heat treatment. At lower heat treatment temperatures (such as 200 °C), the growth of grains is relatively slow, which may be due to the fact that the lower temperature is not enough to provide sufficient thermal energy to promote the diffusion of atoms and the growth of grains. However, the grain size increases significantly when the temperature is annealed to 900 °C, indicating that the atom diffusion rate is accelerated at higher temperatures, thus promoting the coarsening of the grain. The results show that the grain size of Fe-Co-Mo alloy can be controlled by controlling the temperature and time of the heat treatment.

The phase analysis was performed using XRD, as shown in Figure 10. The results show that the phase composition of the sample after 1200 °C hot isostatic pressing, 900 °C annealing, high-temperature solution annealing, and 600 °C precipitation is mainly composed of α-Fe and μ-phases, and no new phase appears in the heat treatment stage. As seen from the magnification in the figure, the diffraction peak of α-Fe at about 52–53 degrees during the solid solution annealing process has obvious migration, and the migration of the diffraction peak returns to normal after the precipitation enhancement, which indicates that the lattice parameters of the material return to a state similar to that after the hot isostatic pressing treatment during the precipitation process. This may be related to the metastable decomposition of the Fe-Co matrix and the formation of B2 order and A2 disorder in Fe-Co matrix. Xie et al. used XRD to analyze the phase composition and crystal structure of Fe-Co-Mo alloys after different heat treatment processes, revealing that the phases present in the alloys include the dominant α-Fe solid solution phase and the diffusely distributed μ-phase [22].

3.3. The Microstructure and Mechanical Properties After Precipitation Hardening

Figure 11 shows the hardness of the material after solution annealing from 1120 °C to 1210 °C at three different HIP treatment temperatures (1200 °C, 1250 °C, 1350 °C). For the 1200 °C HIP treatment, as the solution annealing temperature increases from 1120 °C to 1180 °C, the hardness of the material gradually increases, indicating that in this temperature range, the solid solubility of the alloying element increases, helping to improve the hardness of the material. However, when the temperature rises further to 1210 °C, the hardness begins to decrease, which may be due to the excessive temperature leading to grain coarsening or the dissolution of certain strengthening phases. Microstructural analysis did not reveal any significant grain coarsening, suggesting that the decrease in hardness is more likely associated with the dissolution of certain strengthening phases. In Fe-Co-Mo alloys, the precipitation of nanometer-sized intermetallic μ-phases, with a chemical composition of (Fe,Co)₇Mo₆, plays a critical role in the hardening behavior of the alloy. At higher temperatures, these μ-phases may begin to dissolve back into the matrix, leading to a reduction in hardness. For the 1250 °C HIP treatment, hardness values did not change much throughout the solution annealing temperature range, indicating that the hardness of the material is less sensitive to the solution annealing temperature at higher HIP treatment temperatures. For the 1350 °C HIP treatment, the hardness trend is similar to that of the 1250 °C HIP treatment, and the hardness value does not change much throughout the solution annealing temperature range, further confirming that the hardness of the material is less sensitive to the solution annealing temperature at higher HIP treatment temperatures. For the 1200 °C HIP treatment, the hardness peaks at 1180 °C solution annealing. This may be due to the optimal distribution of solid solubility and the strengthening phase of alloying elements at this temperature, thus providing the highest hardness. The hardness of Fe-Co-Mo alloy shows an obvious change at different solution annealing temperatures, especially for the 1200 °C HIP treatment, as the hardness increases first and then decreases with the increase in temperature. At higher HIP treatment temperatures (1250 °C and 1350 °C), the hardness of the material is insensitive to the change in the solution annealing temperature, indicating that the material may have reached a certain solution saturation state.

In order to further investigate the mechanical properties of Fe-Co-Mo alloys after final heat treatment through solution annealing, the hardness analysis of the sample after precipitation hardening was carried out. Figure 12 shows the hardness of the material after precipitation hardening at different solution treatment temperatures. With the increase in the solution treatment temperature (from 1120 °C to 1210 °C), the sample hardness of all three treatment times changed to different degrees. For HIP samples at 1200 °C and 1250 °C, the hardness increased from 1120 °C to 1210 °C, and it decreased slightly with the further increase in the solution temperature between 1150 °C and 1180 °C. This may be due to the fact that in this temperature range, the solid solubility of the alloying element reaches its maximum, and then the hardness may decrease due to the segregation of the element. The hardness of the HIP treatment at 1350 °C changes gently. It is found that the hardness of the material is about 900 HV after aging strengthening, and the hardness changes gently, indicating that the hardness of the material is less sensitive to temperature changes at higher solution annealing temperatures.

The structure of the precipitated phase of Fe-Co-Mo alloys was further analyzed through transmission electron microscopy. Figure 13 shows a bright field (BF)TEM image of the material after solution annealing at 600 °C, and it shows a selected area electron diffraction (SAED) map obtained from labeled regions A and B, where specific diffraction spots can be observed. According to the SAED pattern shown in Figure 13a,b, the intermetallic compound reinforcement phase particles are identified as the μ-phase. In the bright-field image, μ-phase particles appear as dark regions, and their morphology and distribution can be observed. The electron diffraction pattern is obtained from labeled region B, where the diffraction spots are arranged differently than in Figure 13a. Figure 13a,b show the area of the shaft as [$\overline{1} 1 0$] and [$3 \overline{10} 1$]. In bright field images, the shape and distribution of μ-phase particle A may be different from that of B, and the shape and size of μ-phase particles have an important influence on the strengthening effect of the material. The precipitation of the μ-phase is a key mechanism to improve the hardness and strength of Fe-Co-Mo alloy. Small and evenly distributed particles can usually provide a better strengthening effect. The precipitation enhancement treatment at 600 °C has a significant effect on the state of the μ-phase, thus affecting the final properties of the material. These observations are essential for understanding the strengthening mechanism of the material, optimizing the heat treatment process, and improving the performance of the material. Yuan et al. investigated in depth the effect of Ti(C,N) addition on the microstructure and properties of Fe-Co-Mo carbon-free high-speed steels and observed and analyzed in detail the phase distribution, grain size, and microstructural changes in the alloys due to the addition of Ti(C,N) by using the high-resolution imaging capability of SEM [4].

Figure 14 shows the Vickers hardness changes of Fe-Co-Mo alloy at different holding times (0 to 120 h) and solution annealing temperatures (1120 °C, 1150 °C, 1180 °C, 1210 °C). Resistance to tempering refers to the ability of a material to maintain hardness at high temperatures, which is an important performance indicator for cutting tool materials, such as high-speed steel. The hardness of the samples at all temperatures decreased gradually with the extension of holding time. This may be due to the diffusion of alloying elements at higher temperatures leading to dissolution of the hardening phase or coarsening of the grains. In a certain range, with the increase of solution annealing temperature, the hardness of the material also gradually increased. This suggests that higher annealing temperatures contribute to the formation of a harder solid solution or finer hardened phase particles, thereby increasing hardness. The hardness of 1120 °C solution annealing is the lowest, which may be due to insufficient solution of alloying elements at this temperature, or there are some phase changes that are not conducive to hardness improvement. The hardness after solid solution annealing at 1180 °C and 1210 °C is close to 742.6 HV and 747.6 HV, respectively, indicating that the hardness tends to stabilize with increasing temperature in this temperature range. In the range of 80–120 h, the rate of hardness reduction tends to be gentle. In the case of the 1120 °C solution annealing, it is possible that a peak in hardness could occur at short aging times due to the initial precipitation of hardening phases. When the alloy is first heated to the solution annealing temperature, the solubility of the alloying elements increases, leading to a homogeneous solid solution. Upon short-term aging, these alloying elements may begin to precipitate out as fine, hard particles, which can increase the material’s hardness. Overaging can lead to the dissolution of these precipitates back into the matrix, further reducing the material’s hardness. The presence of a peak in hardness at short aging times is indicative of an optimal balance between the precipitation of hardening phases and their subsequent growth or dissolution. This peak represents a critical point where the material’s resistance to tempering, or red hardness, is maximized. The resistance to tempering of Fe-Co-Mo non-carbon high speed steel is significantly affected by the solution annealing temperature and the holding time. Proper solution annealing temperatures and holding times can optimize the resistance to tempering of the material, thereby improving its cutting performance at high temperatures.

4. Conclusions

(1) During hot isostatic pressing, all three temperatures can achieve complete densification of Fe-Co-Mo carbon-free high-speed steel (relative density of 99.40%), and by controlling the sintering temperature, the morphology and distribution of the µ-phase can be regulated to obtain the Fe-Co-Mo alloy microstructure with a dispersed distribution of reinforcing phases.

(2) Solution annealing can improve the uniformity of the microstructure of Fe-Co-Mo alloy; with the solution annealing temperature increased to 1180 °C, the quantity and particle size of the reinforcing phases presented a tendency to decrease, and the room temperature Vickers hardness decreased to 500.4 HV.

(3) Precipitation hardening can significantly enhance the room temperature Vickers hardness of Fe-Co-Mo alloys, and the hardness of the material can reach about 900 HV after a 600 °C × 3 h precipitation hardening process, which is mainly related to the precipitation of a large number of nano-sized µ-phases in the ⍺-Fe matrix.

(4) By precisely controlling the HIP sintering and subsequent heat treatment processes, Fe-Co-Mo carbon-free high-speed steel with high hardness and excellent resistance to tempering was successfully prepared. After aging treatment, the hardness of the alloy increased significantly to approximately 900 HV, and it possesses good resistance to tempering. This material has great potential for application in high-temperature service environments, especially in fields like aerospace and automotive manufacturing, where there are extremely high demands for the material’s high-temperature performance and wear resistance.