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Investigation into the thermal stability of a novel hot-work die steel 5CrNiMoVNb

  • Zhiqiang Hu EMAIL logo and Kaikun Wang
Published/Copyright: June 10, 2022
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High Temperature Materials and Processes
From the journal High Temperature Materials and Processes Volume 41 Issue 1

Abstract

A novel hot-work die steel 5CrNiMoVNb is developed by optimizing the alloy composition of 5CrNiMoV steel. Thermal stability tests were carried out to compare the hardness evolution of the two steel types. The hardness reduction of 5CrNiMoVNb at 600 and 650°C was only 4.3HRC and 9.6HRC, while that of 5CrNiMoV steel at the same condition was as large as 6.5HRC and 17.5HRC, respectively, which suggests that the thermal stability of the 5CrNiMoVNb steel is more excellent. The thermal stability mechanism of 5CrNiMoVNb was studied based on microstructure analyses and thermodynamic calculations. This suggests that high tempering temperatures cause the coarsening of some carbides and suppress the recovery and recrystallization of the martensite matrix, which is the main reason for the slight decrease in the thermal stability. For the adding of the medium and strong carbide-forming elements, the carbides in 5CrNiMoVNb steel are mainly MC and M23C6 with low coarsening rate coefficient, and the content of these two carbides is almost constant below 670°C. The fine MC and M23C6 carbides showed strong pinning and dragging effects on the dislocations and suppressed martensite recovery and recrystallization. Therefore, the novel hot-work die steel showed excellent tempering softening resistance and thermal stability than 5CrNiMoV steel.

Keywords: 5CrNiMoVNb steel; thermal stability; recovery and recrystallization; carbides; coarsening rate coefficient

1 Introduction

At present, the commonly used hot-work die steels at home and abroad are H13, 3Cr2W8V, 5CrNiMoV, and some new hot-work die steels developed on their basis, which are all typical martensitic hot-work die steels. This type of hot-work die steel is characterized by high contents of medium or strong carbide-forming elements such as Cr, Mo, W, and V. These carbides, such as MC, M2C, M7C3, and M23C6, are controlled to precipitate and achieve the effect of precipitation hardening under reasonable heat treatment processes, thereby improving the thermal resistance, thermal stability, and wear resistance of the steels. Because most hot-work molds are generally exposed to high temperatures and pressures for prolonged periods, it is often prone to thermal fatigue crack, thermal wear, and partial fracture and thus may drastically reduce the actual service life of the die. The most influential parameters on thermal cracking are the surface temperature gradient of the die and the hardness and microstructure of the die material. In some studies, it has been observed that the thermal crack density and crack depth decrease with increasing surface hardness [1]. Thermal wear is a local loss of cohesion or the resulting material loss as a result of the cyclic action of contact stress and internal stress generated by external loads and thermal loads [2,3]. Research shows that die wear, fracture, and other failure modes of hot-work molds during service are ultimately due to the decrease in thermal softening and partial hardness accelerated by higher service temperatures [4,5,6]. Therefore, the softening resistance and thermal stability become the most important service properties for hot-work die steels.

Some investigations into the thermal stability of hot-work die steel suggest that it is dependent on the tempered microstructure and the precipitation strengthening phase [7,8,9]. The tempering temperature mainly influences the size and morphology of martensite and carbides [10]. When the mold is in service at a temperature between 500 and 700°C, the traditional martensitic die steel will exhibit recovery or recrystallization of the microstructure, coarsening of precipitates, and transformation of carbides, which lead to the softening and the reduction in hardness. Under the same tempering temperature and time, the smaller the maximum hardness decreases, the better the thermal stability of the steel [11]. The optimization of alloy composition and adjustment of the heat treatment process are the most commonly used methods to improve the softening resistance and thermal stability of hot-work die steels. Through optimization of the carbide-forming elements such as Cr, Mo, V, W, and Nb, increasing the content of secondary carbides with low coarsening rates and restraining the formation of large-size carbides, such as M23C6 and M3C with high coarsening rates, are conducive to precipitation hardening [3,12,13,14]. In this article, a novel hot-work die steel 5CrNiMoVNb is developed based on the 5CrNiMoV steel. The thermal stability mechanism of the proposed hot-work die steel is analyzed based on the thermal hardness, microstructure evolution, and carbide precipitation.

2 Materials and experiments

2.1 Materials

5CrNiMoV steel is a conventional martensitic hot-work die steel with excellent toughness and hardenability, which is widely used to produce various large or medium forging hammer dies and trimming dies. However, owing to its low temper resistance and thermal stability, 5CrNiMoV steel shows that a short thermal fatigue life and failure fracture easily occurs when it is used under relatively harsh working conditions. Therefore, a new hot-work die steel 5CrNiMoVNb is developed by optimizing the alloy composition. The chemical compositions of 5CrNiMoV and 5CrNiMoVNb steels are shown in Table 1. The 5CrNiMoV steel for the experiment is taken from a large 12-ton hot-work die, and the heat treatment process is oil quenching at 870°C and tempering at 550°C for 2 h. The 5CrNiMoVNb steel is a new type of hot-work die steel developed based on the 5CrNiMoV steel by optimizing the alloy elements through thermodynamic calculations. A 25 kg ingot is obtained by vacuum casting, and hot forging is performed with an eight forging ratio after heating at 1,180°C for 8 h. Afterward, forging is annealed at 850°C, and the next heat treatment process is quenching at 940°C and tempering at 600°C for 2 h. The mechanical properties of the two test steels in a heat-treated state are shown in Table 2. The values of elongation after fracture of 5CrNiMoV and 5CrNiMoVNb are not much different, being 10.13 and 11.40%, respectively. However, the tensile strength, hardness, and impact toughness of the 5CrNiMoVNb steel are increased by 29.4, 32.2, and 33.4%, respectively, compared with those of the 5CrNiMoV steel. This shows that the comprehensive mechanical properties of the 5CrNiMoVNb steel at room temperature are better than those of the 5CrNiMoV steel.

Table 1

Chemical composition of 5CrNiMoV and 5CrNiMoVNb steel

Steel C Si Cr Mn Cu Ni Mo V Nb Fe
5CrNiMoV 0.54 0.25 0.96 0.72 0.12 1.58 0.36 0.074 — Bal.
5CrNiMoVNb 0.55 0.24 0.97 0.70 0.12 1.55 1.80 0.80 0.02 Bal.
Table 2

Mechanical properties of 5CrNiMoV and 5CrNiMoVNb steel

Steel Hardness (HRC) Impact energy (J) Tensile strength (MPa) Elongation (%)
5CrNiMoV 36.2 15.5 1,360 10.16
5CrNiMoVNb 48.3 20.5 1,761 11.40

2.2 Experimental procedures

The initial dimensions of the specimens for the thermal stability tests were 10 mm × 10 mm × 5 mm. To compare the thermal stability of the 5CrNiMoVNb steel with that of the 5CrNiMoV steel, a similar hardness of 42HRC was obtained for the two steels through other two different heat treatment processes, as shown in Table 2. The thermal stability tests were carried out in a box-type resistance furnace. The pretreated specimens were respectively tempered for different times at 600 and 650°C. The hardness of the specimens was measured by an MH-3-type microhardness tester. The microstructures of 5CrNiMoVNb and 5CrNiMoV steel specimens were examined by scanning electron microscopy (SEM; Zeiss Auriga, operated at 20 kV), transmission electron microscopy (TEM; TecnaiF30, operated at 200 kV), and energy-dispersive spectroscopy (EDS). To acquire the SEM specimens, the samples were first ground and polished through multiple passes to remove surface scratches and then corroded by a 4% nitric acid alcohol solution for about 15 s. The TEM specimens were first ground to a thickness of about 50 μm and then round specimens with a diameter of 3 mm were obtained by a punch. Finally, the specimens were thinned by a twin-jet electropolishing device, and the electrolyte was 5% perchloric acid with methanol. The voltage for thinning was about 12 V, and the electrolyte temperature was maintained at −20°C (Table 3).

Table 3

The heat treatment process

Steel The heat treatment process Hardness (HRC)
5CrNiMoV 880°C quenching, 560°C tempering for 2 h 42.4
5CrNiMoVNb 860°C quenching, 620°C tempering for 2 h 41.9

3 Results and discussions

3.1 Analysis of equilibrium phase

As shown in Figure 1, thermodynamic calculations on a multiphase and multicomponent balance were used to analyze the precipitates’ content and their transformation with temperature for 5CrNiMoV and 5CrNiMoVNb steels. For the 5CrNiMoV steel, the main precipitates include M7C3, cementite, M23C6, and MC_ETA when the temperature is below 610°C. The highest mole fraction of the precipitates is carbide M7C3, which is composed of C, Fe, Cr, Mn, Mo, and V. As the temperature increases, the mole fraction of C in M7C3 is almost constant, but the mole fraction of Fe increases a lot by replacing Mn and Cr. When the temperature is higher than 610°C, the carbide M7C3 disappears and the main precipitate transforms into cementite with lower hardness and higher coarsening rate. For the 5CrNiMoVNb steel, when the temperature is lower than 670°C, the main precipitates are MC_ETA, M23C6, and M7C3, and their mole fractions remain almost unchanged. Compared with the 5CrNiMoV steel, the precipitates in the 5CrNiMoVNb steel are significantly increased, especially MC_ETA carbide. As is known that MC has a low coarsening rate and usually exists in the form of fine particles, which contributes to a strong precipitation strengthening effect. The increase of Mo and V elements and the addition of a small amount of Nb element greatly increase the content of MC_ETA, and the content is relatively stable. Thus, the 5CrNiMoVNb steel is likely to have more outstanding toughness and excellent high-temperature thermal stability (Figure 2).

Figure 1 The balanced precipitates of the two steels: (a) 5CrNiMoV and (b) 5CrNiMoVNb.
Figure 1

The balanced precipitates of the two steels: (a) 5CrNiMoV and (b) 5CrNiMoVNb.

Figure 2 The element mole fraction of M7C3 in 5CrNiMoV steel.
Figure 2

The element mole fraction of M7C3 in 5CrNiMoV steel.

3.2 Hardness evolution

As shown in Figure 3, the hardness of 5CrNiMoV and 5CrNiMoVNb steels depends on the tempering time and temperature. As the tempering time increases, the hardness values of the two steels first decrease rapidly and then tend to remain constant. The initial hardness values of the 5CrNiMoV and 5CrNiMoVNb steels are 42.4HRC and 41.9HRC, respectively. After tempering at 600°C for 24 h, the hardness values of the two steels are reduced to 35.9HRC and 37.6HRC, respectively, and the hardness decreases were 6.5HRC and 4.3HRC, respectively. After being tempered at 650°C for 24 h, the hardness values of 5CrNiMoV and 5CrNiMoVNb steels decreased to 24.9HRC and 32.3HRC and the hardness decreases were 17.5HRC and 9.6HRC, respectively. Generally, the hardness reduction of the 5CrNiMoVNb steel at 600 and 650°C is 35 and 45% lower than that of the 5CrNiMoV steel, respectively. This shows that the thermal stability and tempering softening resistance of the 5CrNiMoVNb steel are significantly better than those of the 5CrNiMoV steel. Besides, as shown in Table 4, compared with other commonly used hot-work die steels, such as DM, H21, and H13 steel, the hardness decrease of the 5CrNiMoVNb steel at 650°C for 24 h is at a low level, which suggests that the thermal stability of the 5CrNiMoVNb steel is more excellent.

Figure 3 The dependence of hardness in 5CrNiMoV and 5CrNiMoVNb steels on tempering time and temperature.
Figure 3

The dependence of hardness in 5CrNiMoV and 5CrNiMoVNb steels on tempering time and temperature.

Table 4

The hardness value of several hot-work die steels tempered at 650°C for 24 h

Steel Hardness (HRC) Hardness decrease (HRC)
5CrNiMoV 24.9 17.5
5CrNiMoVNb 32.3 9.6
DM 32.6 14.6
H21 30.4 16.3
H13 25.6 20.9

3.3 Microstructure evolution

The hardness values and the thermal stability usually depend on the tempering microstructure. As shown in Figure 4, the initial microstructure of the two steels prior to the thermal stability test is made up of lath martensite and fine secondary carbides, most of which are evenly dispersed on martensite. Besides the low alloy content of the two steels, there are nearly no large primary carbides. Due to the relatively high tempering temperature for the 5CrNiMoVNb steel during the pretreatment, the initial microstructure is slightly coarsened with recovery characteristics.

Figure 4 The initial microstructure of the two steels by SEM: (a) 5CrNiMoV and (b) 5CrNiMoVNb.
Figure 4

The initial microstructure of the two steels by SEM: (a) 5CrNiMoV and (b) 5CrNiMoVNb.

The tempered microstructure of the two steels at 600°C for 2 and 24 h is shown in Figure 5. At the beginning of the thermal stability test, there was no obvious change in the tempered microstructure of both two steels. After being tempered for 24 h, the carbides are slightly coarsened, and the martensite matrix is recovered in both steel alloys. It is shown that the tempered martensitic lath of the 5CrNiMoV steel is abnormally coarsened and relatively large compared with that of the 5CrNiMoVNb steel. The tempered microstructure of the two steels at 650°C for 2 and 24 h is shown in Figure 6. After tempering for 2 h, the microstructure of both steels is recovered and the martensite lath occurs to disappear. Fine carbides begin to accumulate and grow, and large particles of carbides appear. As the tempering time increases to 24 h, the microstructures of both steels are coarsened and recrystallized, and the martensite lath almost disappears. Some carbides are abnormally coarsened, and most carbides are distributed at the grain boundaries rather than on the matrix. Generally, there are two main factors affecting the thermal stability of the steels. One is the evolution of the second phase caused by hot activation energy and element diffusion, and the other is the recovery and recrystallization of the microstructure due to the release of distortion energy. By comparing the microstructures between 5CrNiMoV and 5CrNiMoVNb steels, it is found that the new hot-work die steel 5CrNiMoVNb has a low carbide coarsening rate and low martensite recovery and recrystallization at the same thermal stability test conditions. Therefore, the thermal stability and the temper softening resistance of the 5CrNiMoVNb steel are significantly better than those of the 5CrNiMoV steel.

Figure 5 The tempered microstructure of the two steels by SEM at 600°C for different tempering times: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, and (d) 24 h-5CrNiMoVNb.
Figure 5

The tempered microstructure of the two steels by SEM at 600°C for different tempering times: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, and (d) 24 h-5CrNiMoVNb.

Figure 6 The tempered microstructure of the two steels by SEM at 650°C for different tempering time: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, and (d) 24 h-5CrNiMoVNb.
Figure 6

The tempered microstructure of the two steels by SEM at 650°C for different tempering time: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, and (d) 24 h-5CrNiMoVNb.

To further analyze the thermal stability mechanism of the novel hot-work die steel 5CrNiMoVNb, the evolution of the microstructure and the types, morphology, and distribution of carbides of the two steels under different thermal stability test conditions are analyzed and compared. The initial microstructure of the two steels by TEM is shown in Figure 7. As shown in Figure 7a, the carbides in the 5CrNiMoV steel are mostly long strips with a width of 30–50 nm. A selected-area diffraction pattern of a long-strip carbide (particle A) shows that the measured angle between the crystal planes (001) and (−110) is 89.53°, which is close to the theoretically calculated value of 90°. After acquiring other crystal face indices through the vector algorithm, the long-strip carbide is calibrated as the orthorhombic M7C3 carbide and the crystal zone axis is (110). The EDS result shows that the long-strip carbide is enriched with Fe, Cr, Mn, and Cu. Compared with M7C3 carbides with a close-packed hexagonal structure, the orthorhombic M7C3 carbides have more excellent thermal stability [15], which shows a positive effect on the thermal stability of the 5CrNiMoV steel. As shown in Figure 7b, most of the carbides in the 5CrNiMoVNb steel are fine and granular. A relatively large carbide (particle B) is analyzed through diffraction (Figure 7d) and energy (Figure 7f) spectra. It suggests that particle B is a face-centered cubic M23C6-type carbide, which is enriched with the elements of Fe, Cr, and Mo. Combining the results of thermodynamic calculations and the analyses of energy spectra of elements, these fine carbides are identified as MC-type carbides. A mass of dislocations is found to be distributed at the boundaries of the martensite slabs, and a few carbides are found around the dislocation cells. This suggests that the movement of dislocations is easily inhibited by carbides and lath boundaries [16,17], resulting in the formation of dislocation cells and high-dislocation-density regions [18,19]. The carbides in the 5CrNiMoVNb steel with an average size of 20 nm are distributed evenly throughout the martensite matrix, which has a stronger pinning effect on dislocations. Therefore, compared with the 5CrNiMoV steel, the 5CrNiMoVNb steel has stronger resistance to temper softening and better thermal stability.

Figure 7 The initial microstructure of the two steels by TEM: (a) 5CrNiMoV, (b) 5CrNiMoVNb, (c) diffraction of particle A, (d) diffraction of particle B, (e) energy spectra of particle A, and (f) energy spectra of particle B.
Figure 7

The initial microstructure of the two steels by TEM: (a) 5CrNiMoV, (b) 5CrNiMoVNb, (c) diffraction of particle A, (d) diffraction of particle B, (e) energy spectra of particle A, and (f) energy spectra of particle B.

Figure 8 shows the tempered microstructures of the two steels by TEM at 600°C for different tempering times. Comparing the martensite and carbides of the 5CrNiMoV steel tempered for 2 and 24 h, the martensite is recovered to a certain extent and clear martensite lath boundaries and a small amount of entangled dislocations can still be observed. Compared with the carbides in the 5CrNiMoV steel tempered for 2 h, the carbides are obviously coarsened, and a small amount of large-particle carbides appears after being tempered for 24 h. For the 5CrNiMoVNb steel, with the extension of the tempering time, the martensite is slightly recovered and the carbides are hardly coarsening. A number of dislocation cells still remain on the martensite matrix. By analyzing the diffraction and energy spectra, particle A in Figure 8d is suggested to be a hexagonal close-packed MC-type carbide with the elements of Mo and V, and the carbides hardly coarsen during the tempering. Research shows that the carbide MC is prone to segregation of Mo atoms to form a Mo-rich layer, which hinders the diffusion of other atoms such as V and Nb into the MC carbide. The Mo-rich layer formed by segregation reduces the coarsening rate by reducing the diffusion rate of C and alloying elements in the MC carbide [20,21]. Besides, element Mo reduces the formation energy of MC carbide and the interface energy between the MC carbide and ferrite [22,23]. Therefore, the (Mo,V)C carbides with a higher degree of coherence with the matrix in the 5CrNiMoVNb steel are fine particles and dispersed on the martensite matrix, which improves the strong plasticity, thermal hardness, and thermal stability of the 5CrNiMoVNb steel to a certain extent.

Figure 8 The tempered microstructure of the two steels by TEM at 600°C for different tempering times: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, (d) 24 h-5CrNiMoVNb, (e) diffraction of particle A, and (f) energy spectra of particle A.
Figure 8

The tempered microstructure of the two steels by TEM at 600°C for different tempering times: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, (d) 24 h-5CrNiMoVNb, (e) diffraction of particle A, and (f) energy spectra of particle A.

Figure 9 shows the tempered microstructures of the two steels by TEM at 650°C for different tempering times. The martensite for the 5CrNiMoV steel obviously recovers and recrystallizes after being tempered at 650 for 2 h. Only a small amount of carbides remains in the matrix, and most of the carbides are coarsened and gathered at the grain boundaries. According to the analysis of the diffraction and the energy spectra for most carbides, the carbide with the highest content is orthorhombic M3C, which has lower hardness and is easy to coarsen. After being tempered at 650 for 2 h, most carbides dissolve into the matrix and a few large-sized carbides are distributed on the grain boundaries of the recrystallized microstructure. Different from the 5CrNiMoV steel, although the martensite is also recovered and recrystallized for the 5CrNiMoVNb steel after being tempered at 650 for 2 h, a large number of MC carbides are still dispersed in the matrix, and M23C6 carbides have not been significantly coarsened. Besides, there are still sparse dislocation lines distributed in the martensite lath, which reflects the pinning and dragging effects of fine carbides on the dislocations. However, as the tempering time increases, a large amount of little carbides dissolve into the matrix and the carbides at the grain boundaries begin to coarsen, which is due to the fact that the fine carbides have a high chemical potential. At the same time, depletion of alloying elements and carbon elements easily occurs inside the grains, which causes the transformation of high coherent carbides to low-coherent and noncoherent carbides. Therefore, the thermal stability and thermal hardness of the 5CrNiMoVNb steel decrease significantly at 650°C, but these are still higher than those of the 5CrNiMoV steel.

Figure 9 The tempered microstructure of the two steels by TEM at 650°C for different tempering times: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, (d) 24 h-5CrNiMoVNb, (e) diffraction of particle A, and (f) energy spectra of particle A.
Figure 9

The tempered microstructure of the two steels by TEM at 650°C for different tempering times: (a) 2 h-5CrNiMoV, (b) 2 h-5CrNiMoVNb, (c) 24 h-5CrNiMoV, (d) 24 h-5CrNiMoVNb, (e) diffraction of particle A, and (f) energy spectra of particle A.

3.4 Coarsening behavior of carbides

Table 5 shows the mass fraction of main carbides in the 5CrNiMoV and 5CrNiMoVNb steel at 600 and 650°C. Based on the TCFE7 and MOBFE3 database of Thermo-Calc software, the application property calculation module is used to calculate the coarsening rate coefficients of the main carbides in the temperature range of 500–725°C, which is shown in Figure 10. For the 5CrNiMoV steel, the main precipitation strengthening phases are 0.281% alloy cementite and 2.720% M7C3 carbide at the temperature of 600°C, where the coarsening rate coefficient of M7C3 is only 0.0247 × 10−29 m3·s−1. When the temperature increases to 650°C, M7C3 disappears and the main precipitation strengthening phases becomes 7.017% alloy cementite and 0.667% M7C3 carbide. At this time, the carbide coarsening rate coefficient of the alloy cementite has increased by two orders of magnitude. The coherent relationship between the large-particle carbide and the matrix is basically lost, and the precipitation strengthening effect is severely weakened. Therefore, the 5CrNiMoV steel shows poor thermal hardness and thermal stability at 650°C. For the 5CrNiMoVNb steel, the precipitation strengthening phases are mainly MC, M23C6, and M7C3 below 670°C, and the content remains almost unchanged as the temperature decreases. The coarsening rate coefficient of MC, M23C6, and M7C3 at 600°C are 9.27 × 10−35, 0.033 × 10−29, and 0.024 × 10−29 m3·s−1, respectively, which suggests that the precipitated phase basically does not coarsen during the process of the long-term thermal stability test at 600°C. When the temperature rises to 650°C, the coarsening rate coefficients of MC, M23C6, and M7C3 increase to 3.66 × 10−33, 0.684 × 10−29, and 0.409 × 10−29 m3·s−1, respectively. Although the coarsening rate coefficients of the three carbides have increased to a certain extent, the MC carbide with the highest content still maintains a very low coarsening rate. Therefore, as the temperature increases, the thermal stability of the 5CrNiMoVNb steel is slightly reduced, but it still has good thermal stability compared with the 5CrNiMoV steel.

Figure 10 The influence of temperature on carbide coarsening rate coefficient.
Figure 10

The influence of temperature on carbide coarsening rate coefficient.

Table 5

Mass fraction of carbides in the two steels at different equilibrium temperatures (%)

M7C3 M23C6 M3C MC
5CrNiMoV steel
 600°C 2.720 0 0.281 0
 650°C 0 0.669 7.017 0
5CrNiMoVNb steel
 600°C 0.848 2.187 0 2.473
 650°C 0.825 2.232 0 2.457

Based on the TEM carbide analysis of the 5CrNiMoVNb steel during the thermal stability test, ImagePro software was used to calculate the carbide sizes of MC and M23C6 with higher contents, as shown in Figure 11. After tempering at 600 and 650°C for 24 h, the size of the MC carbide is almost unchanged. After tempering at 600°C for 24 h, the average diameter of M23C6 carbides is coarsened from 24.84 to 46.31 nm and the coarsening rate is relatively low. When the tempering temperature increases to 650°C, the average diameter of M23C6 carbides is coarsened to 137.44 nm. The estimated coarsening rate coefficient of M23C6 at a temperature of 650°C is about 30 times as much as that at a temperature of 600°C, which is slightly higher than the calculated result. Overall, the precipitation strengthening effect of fine carbides is high and the recovery and recrystallization of the martensite matrix can be suppressed to some extent, for the relatively low coarsening rate of carbides in the 5CrNiMoVNb steel, which is probably the mechanism that the 5CrNiMoV steel possesses excellent thermal stability.

Figure 11 Evolution of main carbide size for 5CrNiMoVNb steel with tempering temperature and tempering time.
Figure 11

Evolution of main carbide size for 5CrNiMoVNb steel with tempering temperature and tempering time.

4 Conclusions

The hardness reduction of the novel hot-work die steel 5CrNiMoVNb at 600 and 650°C is 35 and 45% lower than that of the 5CrNiMoV steel, respectively. The thermal stability and tempering softening resistance of the 5CrNiMoVNb steel are significantly better.

The initial microstructure of the 5CrNiMoVNb steel consists of lath martensite and fine secondary carbides evenly dispersed on the martensite. After being tempered at 600°C and 650°C, the carbides are slightly coarsened and the martensite matrix undergoes recovery and recrystallization, which is the main reason for the decrease of thermal stability.

Due to the reasonable design of the medium and strong carbide-forming elements, the carbides in the 5CrNiMoVNb steel are mainly MC and M23C6 with low coarsening coefficients, and the content of these two carbides is almost constant below 670°C. The fine MC and M23C6 carbides show strong pinning and dragging effects of fine carbides on the dislocations and suppress martensite recovery and recrystallization to a certain extent, which is the thermal stability mechanism of the novel hot-work die steel 5CrNiMoVNb.

Acknowledgments

The authors gratefully acknowledge financial support from the Natural Science Foundation of Suqian City (K202137), Suqian Key Laboratory of High Performance Composite Materials (M202109) and Multifunctional material research and development platform of Suqian University (2021pt04).

  1. Funding information: This work was financially supported by Natural Science Foundation of Suqian City (K202137).

  2. Author contributions: Z.Q.H. designed the experiments, performed the data analysis, and prepared the manuscript. K.K.W. performed the material thermodynamic calculations and revised the manuscript.

  3. Conflict of interest: The authors state no conflict of interest.

  4. Data availability statement: The raw data required to reproduce these findings are available to download from https://data.mendeley.com/drafts/gjy48nwb58.

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Received: 2022-02-06
Revised: 2022-03-25
Accepted: 2022-03-25
Published Online: 2022-06-10

© 2022 Zhiqiang Hu and Kaikun Wang, published by De Gruyter

This work is licensed under the Creative Commons Attribution 4.0 International License.

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  2. Numerical and experimental research on solidification of T2 copper alloy during the twin-roll casting
  3. Discrete probability model-based method for recognition of multicomponent combustible gas explosion hazard sources
  4. Dephosphorization kinetics of high-P-containing reduced iron produced from oolitic hematite ore
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  10. Single-source-precursor synthesis and characterization of SiAlC(O) ceramics from a hyperbranched polyaluminocarbosilane
  11. Carbothermal reduction of red mud for iron extraction and sodium removal
  12. Reduction swelling mechanism of hematite fluxed briquettes
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https://doi.org/10.1515/htmp-2022-0031
Keywords for this article
5CrNiMoVNb steel; thermal stability; recovery and recrystallization; carbides; coarsening rate coefficient
Creative Commons
BY 4.0

Articles in the same Issue

  1. Research Articles
  2. Numerical and experimental research on solidification of T2 copper alloy during the twin-roll casting
  3. Discrete probability model-based method for recognition of multicomponent combustible gas explosion hazard sources
  4. Dephosphorization kinetics of high-P-containing reduced iron produced from oolitic hematite ore
  5. In-phase thermomechanical fatigue studies on P92 steel with different hold time
  6. Effect of the weld parameter strategy on mechanical properties of double-sided laser-welded 2195 Al–Li alloy joints with filler wire
  7. The precipitation behavior of second phase in high titanium microalloyed steels and its effect on microstructure and properties of steel
  8. Development of a huge hybrid 3D-printer based on fused deposition modeling (FDM) incorporated with computer numerical control (CNC) machining for industrial applications
  9. Effect of different welding procedures on microstructure and mechanical property of TA15 titanium alloy joint
  10. Single-source-precursor synthesis and characterization of SiAlC(O) ceramics from a hyperbranched polyaluminocarbosilane
  11. Carbothermal reduction of red mud for iron extraction and sodium removal
  12. Reduction swelling mechanism of hematite fluxed briquettes
  13. Effect of in situ observation of cooling rates on acicular ferrite nucleation
  14. Corrosion behavior of WC–Co coating by plasma transferred arc on EH40 steel in low-temperature
  15. Study on the thermodynamic stability and evolution of inclusions in Al–Ti deoxidized steel
  16. Application on oxidation behavior of metallic copper in fire investigation
  17. Microstructural study of concrete performance after exposure to elevated temperatures via considering C–S–H nanostructure changes
  18. Prediction model of interfacial heat transfer coefficient changing with time and ingot diameter
  19. Design, fabrication, and testing of CVI-SiC/SiC turbine blisk under different load spectrums at elevated temperature
  20. Promoting of metallurgical bonding by ultrasonic insert process in steel–aluminum bimetallic castings
  21. Pre-reduction of carbon-containing pellets of high chromium vanadium–titanium magnetite at different temperatures
  22. Optimization of alkali metals discharge performance of blast furnace slag and its extreme value model
  23. Smelting high purity 55SiCr automobile suspension spring steel with different refractories
  24. Investigation into the thermal stability of a novel hot-work die steel 5CrNiMoVNb
  25. Residual stress relaxation considering microstructure evolution in heat treatment of metallic thin-walled part
  26. Experiments of Ti6Al4V manufactured by low-speed wire cut electrical discharge machining and electrical parameters optimization
  27. Effect of chloride ion concentration on stress corrosion cracking and electrochemical corrosion of high manganese steel
  28. Prediction of oxygen-blowing volume in BOF steelmaking process based on BP neural network and incremental learning
  29. Effect of annealing temperature on the structure and properties of FeCoCrNiMo high-entropy alloy
  30. Study on physical properties of Al2O3-based slags used for the self-propagating high-temperature synthesis (SHS) – metallurgy method
  31. Low-temperature corrosion behavior of laser cladding metal-based alloy coatings on EH40 high-strength steel for icebreaker
  32. Study on thermodynamics and dynamics of top slag modification in O5 automobile sheets
  33. Structure optimization of continuous casting tundish with channel-type induction heating using mathematical modeling
  34. Microstructure and mechanical properties of NbC–Ni cermets prepared by microwave sintering
  35. Spider-based FOPID controller design for temperature control in aluminium extrusion process
  36. Prediction model of BOF end-point P and O contents based on PCA–GA–BP neural network
  37. Study on hydrogen-induced stress corrosion of 7N01-T4 aluminum alloy for railway vehicles
  38. Study on the effect of micro-shrinkage porosity on the ultra-low temperature toughness of ferritic ductile iron
  39. Characterization of surface decarburization and oxidation behavior of Cr–Mo cold heading steel
  40. Effect of post-weld heat treatment on the microstructure and mechanical properties of laser-welded joints of SLM-316 L/rolled-316 L
  41. An investigation on as-cast microstructure and homogenization of nickel base superalloy René 65
  42. Effect of multiple laser re-melting on microstructure and properties of Fe-based coating
  43. Experimental study on the preparation of ferrophosphorus alloy using dephosphorization furnace slag by carbothermic reduction
  44. Research on aging behavior and safe storage life prediction of modified double base propellant
  45. Evaluation of the calorific value of exothermic sleeve material by the adiabatic calorimeter
  46. Thermodynamic calculation of phase equilibria in the Al–Fe–Zn–O system
  47. Effect of rare earth Y on microstructure and texture of oriented silicon steel during hot rolling and cold rolling processes
  48. Effect of ambient temperature on the jet characteristics of a swirl oxygen lance with mixed injection of CO2 + O2
  49. Research on the optimisation of the temperature field distribution of a multi microwave source agent system based on group consistency
  50. The dynamic softening identification and constitutive equation establishment of Ti–6.5Al–2Sn–4Zr–4Mo–1W–0.2Si alloy with initial lamellar microstructure
  51. Experimental investigation on microstructural characterization and mechanical properties of plasma arc welded Inconel 617 plates
  52. Numerical simulation and experimental research on cracking mechanism of twin-roll strip casting
  53. A novel method to control stress distribution and machining-induced deformation for thin-walled metallic parts
  54. Review Article
  55. A study on deep reinforcement learning-based crane scheduling model for uncertainty tasks
  56. Topical Issue on Science and Technology of Solar Energy
  57. Synthesis of alkaline-earth Zintl phosphides MZn2P2 (M = Ca, Sr, Ba) from Sn solutions
  58. Dynamics at crystal/melt interface during solidification of multicrystalline silicon
  59. Boron removal from silicon melt by gas blowing technique
  60. Removal of SiC and Si3N4 inclusions in solar cell Si scraps through slag refining
  61. Electrochemical production of silicon
  62. Electrical properties of zinc nitride and zinc tin nitride semiconductor thin films toward photovoltaic applications
  63. Special Issue on The 4th International Conference on Graphene and Novel Nanomaterials (GNN 2022)
  64. Effect of microstructure on tribocorrosion of FH36 low-temperature steels
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