Effect of vanadium micro-alloying on the microstructure evolution and mechanical properties of 718H pre-hardened mold steel

1. Introduction

Nowadays, the rapid development of plastic mold market has greatly facilitated the production and exploration of pre-hardened mold steels in the industrial field [1], with specific grades of the series AISI P20, 718H (Swedish) and DIN 1.2738 (Germany). Typically, 718H is referred to as pre-hardened mold steel because it is provided after quenching and tempering processes from steelmaking to particular mold manufacturers, and does not require additional heat treatment after the steel is cold machined [2,3]. 718H pre-hardened mold steel offers many advantages in terms of hardness uniformity, polishing properties and machinability, etc [1], and it is provided in tempered sorbite after high temperature tempering treatment. Therefore, the microstructure of 718H steel after quenching must be martensite or lower bainite in order to obtain tempered sorbite. At present, in view of the fierce competition of cost and technology, large section pre-hardened mold steel has been developed with a thickness of over 1000 mm. In order to ensure that the core of the ingot produces martensite or lower bainite microstructure, the 718H steel must have sufficient hardenability.

Micro-alloying is a common method for increasing the hardenability of steel, which indicates that the hardenability is proportional to the alloy content. In recent years, it has been an urgent need to increase the service life of mold steels for environmental and economic reasons. At present, the improvement of mechanical properties of pre-hardened mold steel is mainly achieved by heat treatment [[4], [5], [6], [7]], and the application of micro-alloying is rarely reported.

V is a micro-alloying element commonly used to increase the strength, rather than improve the hardenability of steel. Strength improvement is achieved by precipitation strengthening of precipitates formed during tempering and the combined effect of pinning grain boundaries [8,9]. At the same time, Chen et al. [10] found that V element can effectively improve the hardenability of 40CrNiMoV steel. However, the potential effect of V element on improving hardenability has not been greatly concerned because V mostly exists as carbides in steel. During austenization, V must be dissolved in solid solution instead of forming carbides to improve the hardenability of steel. In addition, a detailed analysis of the different contents of V element on the carbides evolution and impact properties has not yet been carried out. Meanwhile, the mechanism of the V contents on the hardenability is not fully understood.

In this study, the effects of different contents of V (0.1, 0.2, and 0.3 wt%) on microstructure evolution and mechanical properties of 718H pre-hardened mold steel were investigated. The deep characterization of precipitates and microstructures provides a theoretical basis for future industrial application.

2. Experimental

2.1. Materials and heat treatment

Cast ingots with different contents of V (0.1, 0.2, and 0.3 wt%) were smelted by vacuum induction furnace, and the chemical compositions of the test steels are shown in Table 1. In addition, 0.02 wt% RE alloys was added to deoxidize and modify non-metallic inclusion in a vacuum protection atmosphere. Further, cast ingots were forged with a final size of 65 mm × 700 mm × 65 mm (Fig. 1(a)). The heat treatment process included the normalizing, quenching, and tempering. The normalizing condition was 870 °C for 2 h, following by air cooling. The quenching was conducted at 860 °C for 1 h followed by oil cooling. Then, the tempering was performed at 620 °C for 2 h followed by air cooling, as shown in Fig. 1(b).

2.2. Microstructure observation

The original austenite grains were etched by immersion. The etchant used was 6 g picric-acid, 1 g wetting agent, and 100 ml boiling water. The immersion was approximately 40 s while the temperature was kept at 60 °C to avoid crystallization of the picric-acid. Final cleaning was done using alcohol. The dislocation density was analyzed and calculated by X-ray diffraction (XRD) using CuK_α radiation. XRD samples were prepared by using electro-polishing to remove the surface stress. The XRD experiment was performed with 50 kV acceleration voltage and 180 mA beam current, under the 0.02° scanning step from 40° to 105°, and finally using a 2θ/θ scanning mode. The peak intensities were calculated and analyzed by MDI Jade 6.5 software. The dislocation density was estimated based on the method of Williamson-Hall (WH) [11]. Tempered precipitate was detected by XRD and field-emission transmission electron microscope (TEM, FEI Tecnai F20). TEM samples were prepared as described in Ref. [6]. In addition, electron backscatter diffraction (EBSD) technology was used to measure the misorientation gradient of tempered sample. Samples for EBSD analysis were prepared by using ion etching to reduce the surface stress. The EBSD maps were obtained by scanning electron microscopy (SEM) and EBSD HKL system with Channel 5 software. Fracture analysis for impact samples was carried out by SEM.

The chemical composition of the tempered precipitate in 0.2 V steel was determined by using a Cameca local-electrode atom-probe (LEAP) 5000XR. The samples for atom probe tomography (APT) were prepared by two-step electrolytic polishing method. The tip blank with dimensions of 0.5 mm × 0.5 mm × 20 mm was first electrolytically polished with 25 vol.% perchloric acid in acetic acid solution using a direct current voltage of 10-20 V and the final polishing was done with 5 vol.% perchloric acid in butoxyethanol acid solution at 9-15 V [12]. APT test was performed at a voltage pulsing rate at 200 kHz. Image Visualization and Analysis Software (IVAS-3.8.0 version) was used to analyze chemical composition and 3D reconstruction.

Continuous cooling transformation characteristics (CCT) of 0 V, 0.1 V and 0.3 V steels were carried out by an induction heating dilatometric machine (Linseis L78 RITA). Samples with a size of ϕ3 mm × 10 mm were heated to 900 °C at a rate of 2 °C/s and soaked for 900 s, followed by cooling to 25 °C at cooling rates of 0.01, 0.02, 0.05, 0.1, 0.2, 0.5, 1, 2, 5, and 10 °C/s, respectively.

2.3. Mechanical properties tests

The samples with a size of 10 mm × 65 mm × 65 mm were prepared perpendicular to forged direction for hardness testing. Hardness tests were carried out using the Rockwell C scale in a Wilson Rockwell device (LCR-500, LECO Company, St. Joseph, MI, USA), and each sample was measured 10 times to evaluate the cross section uniformity of the material. Cylindrical tensile specimens were prepared along the forging direction with the gage length and diameter of 25 mm and 5 mm, respectively. Tensile tests were carried out at a strain rate of 0.5 mm/min using an AG-100KNG tensile machine (Shimadzu, Kyoto, Japan). The Charpy U-notch specimens with gauge size of 10 mm × 10 mm × 55 mm were prepared with the notch perpendicular and parallel to the forging direction, respectively. The specimens were tested with a 300 J hammer on a pendulum-type impact testing machine (RKP450, Zwick-Roell Company, Ulm, Germany). All mechanical properties were tested at room temperature.

3. Results

3.1. Alloy design

Fig. 2 is the Thermo-Calc calculation result of equilibrium phase diagrams with different V contents according to the chemical compositions (Fe-0.34C-0.31Si-1.53Mn-2.01Cr -0.2Mo-1.0Ni-xV, x = 0.1, 0.2, 0.3 and 0.4). It can be seen that only three types of precipitates M₇C₃, M₂₃C₆ and MC were found in 0.1 V, 0.2 V, and 0.3 V steels. However, when the V content was increased to 0.4 wt%, only two types of precipitates M₇C₃, M₂₃C₆ appeared. In order to increase the strength and hardness, MC precipitates need to be generated during the normal tempering temperature (500-650 °C [6]). Therefore, the content of V element should be less than 0.3 wt% in this study.

3.2. Microstructural characterization

Fig. 3 indicates the XRD spectra of tempered samples with different V contents. The results showed that austenite peaks did not appear in all samples. It shows that the retained austeniteis very low due to the volume fraction of retained austenite (V_RA) is below the detection limit [13]. However, α (martensite) peaks are clearly visible in all samples. In addition, the dislocation density is estimated to be 3.54 × 10¹⁴, 3.66 × 10¹⁴, 3.83 × 10¹⁴, and 3.77 × 10_14*m^-2 corresponding to 0 V, 0.1 V, 0.2 V, and 0.3 V steels. The result shows that there is no significant change in the dislocation density with different V elements.

The CCT diagrams of 0 V, 0.1 V and 0.3 V steels were presented in Fig. 4. It is obvious that all CCT diagrams exhibit the characteristics of complex microstructures and continuous phase transformation as the cooling rate increases from 0.01 to 10 °C/s [6]. The shapes of the CCT diagrams of test steels with different V contents are similar. Pearlite (P), bainite (B), and martensite (M) transformation occur in sequence. The difference between these diagrams is that the P transformation ranges of 0.1 V and 0.3 V steels are shifted to a shorter periodic side compared with 0 V steel. The detailed result is that the P transformation took place at a cooling rate of 0.05 °C/s in V-free steel, while it did not occur for 0.1 V and 0.3 V steels (Fig. 4(b, c)).

Fig. 5 shows that optical metallographs (OM) of original austenite grain structure of samples after quenching with different V contents. The results show that the samples of 0 V and 0.1 V steels exhibit similar grain structure and grain size after quenching. However, when the V content is increased to 0.2 wt%, the mixed grain structure and zigzag grain boundary appeared, especially in 0.3 V steel (Fig. 5(c, d)). Fig. 6 indicates EBSD analyses of samples tempered at 620 °C with different V contents, and high-angle grain boundaries (HAGBs) with misorientations of ≥15° are expressed by black lines (Fig. 6(e-h)). In addition, HAGBs were determined as the effective grain boundary and its size was accurately measured by Image Pro Plus 6.0 software (IPP 6.0). The results are summarized in Table 2, and the effective grain size was 1.126, 1.085, 1.052, and 1.035 μm corresponding to 0 V, 0.1 V, 0.2 V, and 0.3 V steels, respectively.

Fig. 7(a, b) indicates the XRD pattern of quenched precipitates with different V contents. The results show that no XRD peaks of precipitates are found in 0 V steel. However, MC precipitates appeared when the V element was added. In addition, MC, M₃C and M₇C₃ were detected in V-added steels when tempered at 620 °C. For 0 V steel, only M₃C and M₇C₃ were detected [6]. Simultaneously, the structure of tempered precipitates with different V contents was determined by TEM with SAED pattern, as shown in Fig. 8. The striped particles in Fig. 8(a, b) are identified as M₃C carbide (zone A and C). The particles in Fig. 8(a, c) is the M₇C₃ carbides with hexagonal close packed (hcp) structure (zone B and D). In addition, large numbers of fine scale particles with a size from 5 to 10 nm were detected in V-added steels (Fig. 8(b-d)). However, no this type of particles were found in 0 V steel (Fig. 8(a)). In addition, fine scale particles can pin the dislocation and restrain dislocation movement in 0.1 V, 0.2 V, and 0.3 V steels (Fig. 9), which is consistent with our previous studies [14].

The atomic distribution of different elements of fine scale particles in 0.2 V steel was characterized by APT technique, as shown in Fig. 10. It shows that the C, Mo, and V atoms were segregated in the same position, while the Ni, Mn, Cr and Fe atoms were uniformly distributed in matrix (Fig. 10(a)). In addition, a box with a certain volume was selected and the concentration distribution was obtained in Fig. 10(b, c). In this study, only the C, Mo, Fe and V elements were selected to study the elemental distribution of tempered precipitates. The results show that the C, Mo, and V elements are segregated at the same position of the inner layer of precipitate in zone B; the outer layer is enriched with Fe atoms (Fig. 11). The precipitates in zones A and B have a similar chemical composition and their sizes are less than 20 nm (Fig. 10(b, c)). Combined with the XRD results in Fig. 7(c) and the description in Ref. [8,14], the precipitates in zones A and B are identified as MC precipitate. Fig. 12 indicates the distribution of atoms of different elements of tempered precipitate in 0 V steel by APT technique. It is found that the precipitate is mainly enriched in elements such as C, Cr, Fe, Mn, Mo, and evenly distributed, but the size is significantly larger than MC, which is consistent with the characteristics of precipitates M₃C and M₇C₃.

3.3. Mechanical properties

The hardness was measured with different V elements of test steels and the results are shown in Fig. 13(a). It indicates that the hardness is 33.51, 34.18, and 34.86 HRC corresponding to 0.1 V, 0.2 V, and 0.3 V steels, respectively, higher than 0 V steel (29.24 HRC).

Fig. 13(b) displays the yield strength (YS), tensile strength (UTS), section shrinkage, and elongation for the steels with different V contents. The addition of V element led to a remarkable increase in YS and UTS, but a decrease in total elongation. The YS is increased from 855 MPa (0 V) to 967, 990, 997 MPa corresponding to 0.1 V, 0.2 V, and 0.3 V steels. In contrast to the increase in YS and UTS, the elongation decreased from 67% to 62%, 62%, 61%, respectively.

The Charpy U-notch absorbed energy vs different V contents of the steels is shown in Fig. 13(c). The impact energy decreases gradually with the V addition. The transverse impact energy is decreased from 90.5 J to 78.4, 65.8, 57.5 J corresponding to 0 V, 0.1 V, 0.2 V, and 0.3 V steels, respectively. Similar phenomenon also occurred in the longitudinal impact energy values of the test steels.

Fracture analysis of impact samples of test steels were carried out to reveal the fracture morphologies at low and high magnifications, respectively, as shown in Fig. 14. Obviously, the overall fracture morphology and the zone B for all samples were similar. Various sizes of dimples were observed by high magnification, which revealed a ductile fracture. In addition, the strip and spherical inclusions were well distributed in the dimples of test steels. EDS analysis indicated that the inclusions were RE-oxy-sulfides and Mn-S inclusions (Fig. 14(e)), which is consistent with our previous studies [14].

4. Discussion

4.1. Microstructural evolution

Compared with 0 V steel, the dislocation density of the test steels with different V contents did not change significantly, even if the fine scale precipitates can pin the dislocation (Fig. 9). However, the dislocation does not slip due to non-deformation, resulting in less increment of static dislocation density. In addition, Fig. 4 displays that the synergistic effect of V elements reduces the P transformation range compared with 0 V steel. The main reason is that there is a large lattice mismatch between V and Fe atoms, resulting in a decrease in the solubility limit of V atoms in steel [15]. It encourages the segregation of dissolved V atoms at the grain boundary to reduce the grain boundary energy and inhibit the diffusion of C atoms [10]. As a result, it delays the nucleation of ferrite/pearlite and eventually improves hardenability of V-added steels. These indicate that if most of the V atoms in steels combine with C atoms to form carbides, the hardenability is deteriorated during the subsequent treatment process. Therefore, the dissolved V content in austenite determines the hardenability. The dissolved V content was calculated at 860 °C (quenching temperature) by using the solid solubility formula and ideal stoichiometry of C and V atoms in VC carbide. The solid solubility formula is shown in Eq. (1) [16]:

Log([V]·[C])=6.72-$\frac{9500}{T}$ (1)

where [V] and [C] are the mass fractions of solution elements in test steel (wt%) and T is temperature (K). Hence, the dissolved V and C elements in austenite are presented in Eq. (2):

[V]_γ·[C]_γ=0.0216 (2)

Simultaneously, the ratio of V/C should maintain the ideal stoichiometry as in Eq. (3):

V_vc/C_vc=4.24 (3)

Therefore,

$\frac{0.1-[V]_γ}{0.34-[C]_γ}$=4.24 (4)

The dissolved V and C elements can be obtained using Eq. (2) and (4):

[V]_γ=0.0651,[C]_γ=0.3317

As a result, the dissolved V in austenite for 0.1 V steel was 0.0651 wt%, accounting for 65.1% of the total weight (0.1 wt%). For 0.3 V steel, [V]_γ was 0.0753 wt%, accounting for only 25.1% (0.3 wt%). It suggests that most of the V and C atoms combined to form VC carbides in 0.3 V steel. Therefore, the hardenability of both 0.1 V and 0.3 V steels was improved by the dissolved V elements. However, compared with 0.1 V and 0.3 V steels, the ability to improve the hardenability is decreased due to the formation of more VC carbides.

Previous studies have revealed that the grain size decreased with different V elements in steels [8]. The mechanism of grain refinement with different V contents may be related to the delay of recrystallization during the hot deformation process. In this study, the effective grain size of the test steels decreases continuously with addition of V element (Table 2). This can be explained by the fact that the V element will refine the martensite packets and blocks of the test steels. Smaller martensite size shows higher amounts of HAGBs between packets [17]. Simultaneously, mixed grain structure and zigzag grain boundary appear of samples after quenching in 0.2 V and 0.3 V steels (Fig. 5(c, d)). The main reason can be attributed to the uneven distribution of MC carbides and the dissolved V elements in steel during the austenization. It can be concluded from the Fig. 7(b) that MC precipitates appear in the quenched samples in V-added steels. However, no precipitates were found in the quenched samples of V-free steel. In addition, there was a report that the fine scale precipitates had a pinning effect on grain growth with the size less than 50 nm [18]. In this study, the sizes of MC precipitates were less than 50 nm in V-added steels (Fig. 8(b-d)). During the grain growth, the fine scale MC precipitates and the dissolved V elements in steel would impede the movement of grain boundary. In addition, the same heat treatment process was selected in order to ensure the accurate comparison of mechanical properties. For 0 V and 0.1 V steel, the selected austenization temperature (870 °C) is about 30 °C higher than its Ac₃ temperature (829, 840 °C, respectively, in Fig. 4), resulting in a uniform grain size. But for 0.3 V steel, the austenization temperature is only 14 °C higher than its Ac₃ temperature (856 °C). The incomplete austenization process may also be one of the reasons for the occurrence of mixed grain structure and zigzag grain boundary in 0.2 V and 0.3 V steels. These explanations are consistent with experimental results in which the mixed grain structure is more pronounced with the increase of V contents.

The difference between the precipitation sequences of V-free and V-added steel is the presence of MC. Microalloying of 0.1 to 0.3 wt% V elements is confirmed to be the main reason of MC precipitation according to XRD characterization (Fig. 7) and Thermo-Calc calculations (Fig. 2). V atoms will force the C atoms to diffuse themselves and form the VC carbides (Fig. 11) [14]. Simultaneously, V atoms seem to have the ability to decrease the diffusion rate of C and other alloy atoms by forming strong carbides, resulting in preventing the coarsening of precipitates [9].

4.2. Mechanical properties

The contribution of various strengthening mechanisms to the yield strength of test steels is derived from following Eq. (5) [19,20]:

σ_YS=Δσ₀+Δσ_ss+Δσ_GB+$\sqrt{(Δσ_{DIS}^2+Δσ_{Prec}^2)}$ (5)

where σ_YS is yield strength (MPa); Δσ₀ is intrinsic strength (MPa) and is identified as about 85 MPa [21]; Δσss is solid solution strengthening (MPa). The contributions of the yield strength were estimated by reference and experimental results. Eq. (6) represents the solid solution strengthening of different components [19,21]:

Δσ_ss=32.34[Mn] +83.16[Si]+360[C]+33[Ni]+11[Mo]+354.2[N] (6)

where [Mn], [Si], [C], [Ni], [Mo], and [N] are the mass fraction of solution elements, respectively (wt%). The contribution of C and carbide forming elements (Cr) is considered negligible due to the low solubility during high temperature tempering [21,22]. The contribution of solid solution strengthening is estimated to be approximately 113 MPa corresponding to 0 V, 0.1 V, 0.2 V and 0.3 V steels, respectively based on Eq. (6).

Δσ_GB corresponds to the grain refinement strengthening. The following Eq. (7) represents the grain refinement strengthening:

Δσ_GB=K_HP·d^-1/2 (7)

Zhang et al. [23] reported that the martensitic block was the minimum substructure unit in controlling the strength. Daigne et al. [24] indicated that the grain boundary effect is made up of the size of packet or lath width. Shibata et al. [25] believed that the improvement of strength due to the existence of block boundaries. The martensitic block was successfully indicated with an assistance of EBSD for SEM and a definition of grain misorientation higher than 15° as the block boundary [26]. Chen et al. [27] also proposed that the grain boundaries were defined where the misorientation exceeds 15° when calculating the strengthening mechanism of low carbon 3Mn-1.5Ni steel. Therefore, the Hall-petch term (d) should be replaced by the block boundary strengthening in Eq. (7). Here, d is the effective grain size, which is summarized in Table 2. In addition, the value of KHP is considered to be 0.19 MPa m^-1/2 according to Kim et al. [22] study on the strengthening mechanism of tempered martensite steel. It is complies with the value of KHP is 0.21 MPa m^-1/2 of the block boundaries reported by Shibataet et al. [25] and 0.20 MPa m^-1/2 reported by Wang et al. [28]. The contribution of grain refinement strengthening is estimated to be about 179.05, 182.40, 185.24, and 186.75 MPa corresponding to 0 V, 0.1 V, 0.2 V and 0.3 V steels, respectively.

Δσ_DIS is dislocation strengthening that can be calculated using the following Eq. (8):

Δσ_DIS=αMGbρ^1/2 (8)

where M is 3.0 (Taylor factor), α is 0.4 (constant), G is 78.5 GPa (shear modulus), b is 0.248 nm (Burgers vector). ρ represents dislocation density of test steels (Fig. 3). The contribution of dislocation strengthening is estimated to be about 439.5, 446.93, 457.19 and 453.60 MPa corresponding to 0 V, 0.1 V, 0.2 V and 0.3 V steels, respectively based on Eq. (8).

In addition, the volume fraction and morphology of precipitates plays an important role in improving the yield strength of test steels [5]. Eq. (9) represents the precipitation strengthening according to Orowan relationship [29]:

Δσ_Prec=$\frac{0.538Gbf^{1/2}}{D} ln\frac{D}{2b}$ (9)

where Δσ_Prec represents the precipitation strengthening. D and f represent the diameter and volume fraction of the precipitates, respectively. The following Eq. (10) is used to calculate the f [30]:

f=V_c·ρc (10)

where V_c represents the volume of precipitates; ρc represents the number of precipitates per unit volume of matrix. Many literatures reported that the precipitation strengthening was calculated using TEM frames [30]. However, the accuracy of the calculated results is low due to the statistical error of the thickness of TEM thin foil thickness and the number of precipitates. In our study, the increment of precipitation strengthening can be calculated by subtracting of Δσ₀, Δσ_ss, Δσ_GB, and Δσ_DIS from total yield strength. Based on Eq. (5), the contribution of precipitation strengthening is estimated to be about 187.81, 379.93, 398.91 and 411.21 MPa corresponding to 0 V, 0.1 V, 0.2 V and 0.3 V steels, respectively. Therefore, it indicates that the MC precipitation strengthening is the main contribution strengthening to V-added steels according to the calculation of various strengthening mechanisms.

The decrease in elongation can be attributed to the increase of the strength values [31]. However, the cause of the reduction in impact energy is more complicated. Impact energy of martensitic steel can be increased by smaller austenite grain size and finer block boundaries [14]. Studies showed that the direction of crack could change as the cracks attempt to propagate across HAGBs, thus delaying the crack propagation during the test of impact [32]. But the study shows that the effective grain size decreases continuously of test steels with different V contents, which excludes the main effects of HAGBs on the decrease in impact energy. In addition, the modification of non-metallic inclusion and purification of molten steel by RE alloys also play significant roles in improving impact toughness [33]. In our study, there have been fewer changes in these two areas.

Based on the dislocation theory and the Hall-Petch formula, Cottrell [34] proposed the expression for the extended critical stress of crack cleavage fracture (Eq. (11)):

σ_f=$\frac{2Gγ_m}{ ky}$·$d^{-\frac{1}{2}}$ (11)

where σ_f is the critical stress of crack propagation; G is shear modulus; γ_m is the plastic work of crack propagation; k_y is the term of yield constant; d is effective grain size, and the data is shown in Table 2.

According to Eq. (11), σ_f is proportional to γ_m and inversely proportional to d of the material. However, the effective grain size decreased only from 1.126 μm to 1.035 μm with a small amount of variation. Therefore, the cause of the decrease in impact toughness can be attributed to the interaction between precipitates and dislocations. Recent literatures reported that the precipitates tend to heterogeneous nucleate on dislocations in order to decrease the strain field due to the difference in lattice parameter between ferrite and precipitates [35,36]. Accordingly, fine scale MC precipitates can pin dislocations, which lead to decrease in its ability to move. A large number of immovable dislocations will increase the tendency of dislocation accumulation and internal stress, making plastic deformation difficult to perform, thereby reducing γm [37]. This will also lead to a decrease in σ_f, which increases the tendency of brittle cracking. In addition, Najafi et al. [31] also showed that the fine scale precipitates can lock dislocations and make micro-alloying samples vulnerable to instant loading, resulting in a gradual decrease in impact toughness.

5. Conclusions

The effect of different contents of V (0.1, 0.2, and 0.3 wt%) on microstructure evolution and mechanical properties of 718H steel were investigated. The conclusions can be drawn as follows.

(1) Pearlite transformation ranges of 0.1 V and 0.3 V steels are shifted to a shorter periodic side compared with 0 V steel. The main reason is that the segregation of dissolved V atoms decreases the grain boundary energy and inhibits the diffusion of the C atoms, eventually delays the nucleation of ferrite/pearlite. However, compared with 0.1 V and 0.3 V steels, the ability to improve the hardenability is decreased due to the generation of more VC carbides.

(2) The uneven distribution of MC carbides and the dissolved V elements during the austenization as well as the incomplete austenization process are the reasons for the occurrence of mixed grain structure and zigzag grain boundary in 0.2 V and 0.3 V steels.

(3) The YS is increased from 855 MPa (0 V) to 967, 990, 997 MPa corresponding to 0.1 V, 0.2 V, and 0.3 V steels. It indicates that the MC precipitation strengthening mainly contributes to the V-added steel by analyzing various strengthening mechanisms.

(4) The transverse impact energy is decreased from 90.5 J to 78.4, 65.8, 57.5 J corresponding to 0 V, 0.1 V, 0.2 V, and 0.3 V steels, respectively. Similar phenomenon also occurred in the longitudinal impact energy values. A large number of immovable dislocations pinned by fine scale MC precipitates will increase the dislocation accumulation, internal stress and brittle cracking, resulting in a gradual decrease in impact toughness with the V addition.

Acknowledgements

The work was financially supported by the National Key Research and Development Program of China (No. 2016YFB0300401). This work also was supported by Cooperation Program of Hubei province and Chinese Academy of Sciences (The Research and Development of Key Technologies for Special Steel of Homogeneous High Performance).

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