Effect of Nb Content on the Behavior of Primary Carbides in 0.4C-5Cr-1.2Mo-1V Steel

Abstract

Niobium (Nb) microalloying can improve the material properties of H13 steel (0.4C-5Cr-1.2Mo-1V steel), but it also affects the natures of the primary carbides. Therefore, the effect of Nb content and cooling rate on the behavior of primary carbides in H13 steel was studied in this paper. The matrix structure was obtained by chemical etching, and then the formation location of primary carbides was identified by electron probe microanalysis (EPMA). The three-dimensional (3D) characteristics, including morphology, number density, and size, were obtained by a non-aqueous electrolysis method. The enrichment of alloying elements in the last-to-solidify region leaded to the formation of primary carbides during the solidification. The Ti4C2S2 phase precipitated first, and then the Mo-Cr-rich carbide was formed around the Ti4C2S2 phase. During the cooling process, the Ti4C2S2 phase partly transformed into Nb-rich carbide and then further partly transformed into V-rich carbide. There is a huge difference between the two-dimensional and three-dimensional morphologies of the primary carbides. As the Nb content increased, the size of last-to-solidify region decreased gradually and the size and number density of primary carbides in the 3D observation increasingly increased. However, as the decrease of the cooling rate, the size of primary carbides increased rapidly and the number density of primary carbides decreased markedly. The thermodynamic and kinetics calculation results agreed well with the experimental observations.

1. Introduction

H13 steel (0.4C-5Cr-1.2Mo-1V steel) is the most widely used hot work die steel and used in hot rolling, hot extrusion and hot forging because of its combination of notable hot strength and remarkable toughness, especially the thermal fatigue resistance. H13 steel exhibits excellent mechanical properties after being quenched at 1050°C and tempered twice at 600°C.¹⁾ To further improve the performance of H13 steel, laser surface hardening,²⁾ deep cryogenic treatment,³⁾ and Nb microalloying⁴⁾ have also been employed. Nb microalloying has been widely used in refining grain size and improving strength and impact toughness of steels.^5,6) The total substitution of vanadium in H13 steel by 0.07 wt.% niobium can effectively suppress grain growth and coarsening MC carbide.⁷⁾ According to Elias’s results, the partial substitution of vanadium by niobium does not lead any significant change in the hardness of the H13 hot work tool steel, on the contrary, the new steel has a smaller austenite grain size and a finer second carbide size distribution, indicating a higher potential toughness.^8,9) Maloney found that the addition of 0.035 wt.% niobium can significantly improve the high temperature toughness and heat check resistance of H13 steel without any decrease in hot hardness properties.¹⁰⁾

However, the addition or substitution of niobium will inevitably change the characteristics of the primary carbides in H13 steel. Large amounts of primary carbide containing elemental Nb can be found in the commercial Nb microalloying H13 steel.¹¹⁾ On the one hand, the primary carbide cannot be eliminated completely during the heat treatment because of the high thermal stability, especially the primary carbide containing elemental Nb.^12,13) On the other hand, the precipitation of primary carbides reduces the solid solution quantity of elemental Nb in the matrix, leading to a decrease in the content of secondary carbides precipitated during the tempering process. The large primary carbides are deleterious to the impact toughness of H13 steel and even act as the crack sources to reduce the service life of H13 steel.^14,15) At present, few people are involved in studying the effect of Nb content on the behavior of primary carbide in H13 steel, especially from the perspective of the three-dimensional natures of primary carbide. It is hoped that the research work can provide a guiding role for the effective control of primary carbides in H13 steel.

2. Experimental Section

Experimental H13 ingots with different amounts of elemental Nb were manufactured in a vacuum induction furnace. First, pure iron (99.99%) was melted in a MgO crucible, and then Si (99.5%), Mn (99.25%), Cr (99.69%), Mo (99.99%), Nb (99.99%) and V-Fe (51.84% V and 46.56% Fe) were added into the MgO crucible. Secondly, the temperature of the molten steel was kept at 1600°C for 5 min to homogenize its composition. Finally, the molten steel was casted into a cast iron model. The obtained H13 ingot was stripped and air cooled to room temperature after 10 min. The chemical composition of H13 ingots was determined at the National Analysis Center for Iron and Steel, and the results are shown in Table 1. The Nb-0, Nb-1, Nb-2, and Nb-3 samples were cooled in air, and the Nb-4 sample was cooled in vacuum induction furnace. The cooling rate in furnace is much lower than in air.

Table 1. Chemical composition of samples, wt.%.

Samples	C	Si	Mn	P	S	Cr	Mo	V	Ti	Nb	Cooling condition
Nb-0	0.38	1.00	0.56	0.013	0.008	4.98	1.32	0.99	0.011	0	Cooled in air
Nb-1	0.41	1.00	0.58	0.013	0.007	5.30	1.38	0.96	0.014	0.02	Cooled in air
Nb-2	0.42	1.01	0.6	0.013	0.007	5.36	1.35	0.95	0.015	0.04	Cooled in air
Nb-3	0.42	1.00	0.58	0.013	0.007	5.22	1.32	0.98	0.014	0.08	Cooled in air
Nb-4	0.40	1.00	0.57	0.013	0.007	5.30	1.34	0.98	0.015	0.08	Cooled in furnace

The solidification microstructure of H13 steel ingot was obtained after being corroded for 5 min by the hydrochloric acid water solution with volume ratio 1:1 at 70–80°C. Two samples for each H13 ingot were taken on both sides of the centerline at 30 mm from the bottom, and the samples were located in the equiaxed crystal region. The automatic inclusion analysis system (EVO18-INCA steel) was used to analyze the chemical composition of the primary carbide after being polished. The statistical area was 5 mm×5 mm, and the minimum size of inclusions was 1 μm. The morphology and type of the inclusions were analyzed through scanning electron microscope (SEM) (FEI Quanta-250) equipped with an energy dispersive spectroscope (EDS) (XFlash 5030). A non-aqueous electrolysis method was used to investigate the 3D natures of primary carbides in H13 steel ingots. Non-metallic inclusions were partially extracted from the sample matrix, and the extracted area was 10 mm×10 mm. The electrolyte was composed of 1% tetramethylammonium chloride, 10% acetylacetone, and 89% methanol (value fraction). The electrolysis voltage was 20 V, and the electrolysis time was 300 s. At least 150 primary carbides were observed in each sample by 3D observation with SEM to determine their number density and size. To obtain the precipitation mechanism of the primary carbide during solidification, the matrix microstructure of samples was investigated with an optical microscope (OM) (DM4M) after being etched with an alcohol solution containing 4% nitric acid (volume fraction). We found the reticular last-to-solidify region in the matrix microstructure at 200 magnification times. ImageJ software was used to measure the size of last-to-solidify region with a mean linear intercept method, including a small amount of last-to-solidify region that is not closed. We first got the length of the measuring line, and then get the number of intersections between the measuring line and the last-to-solidify region. The ratio of the length to the number of intersections was the average size of last-to-solidify region. Each sample was measured at least ten times. The electron probe microanalysis (EPMA) (EPMA-1720) was used to analyze the chemical composition of the last-to-solidify region and primary carbides. Finally, commercial thermodynamic software Thermo-Calc 2017b, including Scheil and equilibrium models, was used to calculate the precipitation mechanism of primary carbides and the influence mechanism of Nb content on primary carbides in H13 ingots. The total system was set as 1 g. The Scheil model assumes that the alloying elements are homogeneous in liquid and that diffusion did not occur in the solid steel.

3. Results

3.1. Matrix Structure

The matrix structure is shown on the left of Fig. 1. The sampling location is in the equiaxed crystal region. The bright white area observed by OM is the last-to-solidify region, and primary carbides were all found in the last-to-solidify region. The 3D morphology of primary carbide is shown on the right of Fig. 1. The 3D morphology of primary carbide is a typical dendritic structure, and their size reaches 60 μm. Compared with the 2D morphology of primary carbide obtained by other researchers, it can be seen that the natures of the primary carbide is one-sided in the 2D observation, which agrees well with our previous work.¹⁶⁾ Therefore, it is of great significance to study the characteristics of primary carbides in the 3D observation.

Fig. 1.

Solidification microstructure of H13 ingot and sampling location. On the left is the precipitation position of primary carbide obtained by chemical etching, and on the right is the 3D morphology of primary carbide obtained by non-aqueous electrolysis method. (Online version in color.)

The morphologies of the last-to-solidify region are shown in Fig. 2. The last-to-solidify region of Nb-0, Nb-1, Nb-2, and Nb-3 samples is a typical network structure. However, the network structure of the last-to-solidify region in Nb-4 sample is been broken, indicating that the cooling rate has a great influence on the morphology of last-to-solidify region. The chemical composition of last-to-solidify region was further analyzed by EPMA. Taking the Nb-1 sample as an example, and the results are shown in Fig. 3. Elemental Cr, Mo, V, and C are largely enriched in the last-to-solidify region, and large amounts of primary carbides, consisting of elemental Cr, Mo, V, C, and/or Nb, were only observed in the last-to-solidify region, as shown the red dot in Fig. 3. The red dot means the mass fraction of corresponding element in this area is the highest. Therefore, the red dot is the primary carbide. During the growth of the grain in the equiaxed crystal region, elemental Cr, Mo, V, C, and Nb are rejected to the liquid and may exist in relatively high concentrations in the final solid that forms along the solidification boundaries, which has been identified widely by other researchers in steel and Ni-base superalloys.^{17,18,19,20,21)} Therefore, the last-to-solidify region size is the original austenite grain size.

Fig. 2.

Morphologies of last-to-solidify region in samples, (a) Nb-0 sample, (b) Nb-1 sample, (c) Nb-2 sample, (d) Nb-3 sample, (e) Nb-4 sample.

Fig. 3.

Chemical composition of last-to-solidify region in Nb-1 sample analyzed by EPMA. (Online version in color.)

The variation of Nb content with last-to-solidify region size is shown in Fig. 4. The error bars represent the standard deviation. Since the network structure of last-to-solidify region in the Nb-4 sample is broken, making their size difficult to be counted, the last-to-solidify region size of the Nb-4 sample is not shown in Fig. 4. As the Nb content increases, the last-to-solidify region size gradually decreases from 85.7 to 64.9 μm. Therefore, the addition of elemental Nb refines the solidification structure of H13 ingots.

Fig. 4.

Variation of Nb content with last-to-solidify region size.

3.2. Two-dimensional Characteristics of Primary Carbide

The 2D morphologies and types of inclusions observed in all samples are the same, and the specific morphology and composition are shown in Fig. 5. The typical primary carbides were divided into three types: 1, V-rich, as shown the gray area in Figs. 5(a)–5(e); 2, Mo-Cr-rich, as shown the white area in Figs. 5(b)–5(e); 3, Nb-rich, as shown in Fig. 5(f). Moreover, certain amounts of MnS inclusions were found, as shown the black area in Figs. 5(c), 5(e). The shape of primary carbides is strip, and their sizes were approximately 10 μm. The MnS inclusion was present either alone or combined with primary carbides. The automatic inclusion analysis system was used to analyze the chemical composition of primary carbides containing elemental Nb, and the results are shown in Fig. 6. As the Nb content increases, the Nb content in the primary carbide gradually increases as well.

Fig. 5.

2D morphologies of the primary carbide. (Online version in color.)

Fig. 6.

Composition of primary carbide containing elemental Nb in samples. (Online version in color.)

3.3. Three-dimensional Characteristics of Primary Carbide

Since the number density and size of the primary carbide in 2D observation have larger errors with actual results, we obtained the characteristics of primary carbide in the 3D observation.¹⁶⁾ Figure 7 shows the 3D morphologies of primary carbides in samples. The primary carbides are all distributed along the grain boundaries for both cooled in air samples (Nb-0, Nb-1, Nb-2, and Nb-3) and cooled in furnace sample (Nb-4). However, the network distribution of primary carbides was interrupted when the cooling condition was changed from in air to in furnace, which is consistent with the results in Fig. 2. The 3D morphology of primary carbides is a typical dendritic structure, and the detail distribution can be found in our previous work.¹⁶⁾ Nb content and cooling rate have little influence on the 3D morphology of primary carbides.

Fig. 7.

3D morphologies of primary carbide in all samples. (Online version in color.)

The composition of primary carbides in all samples were basically the same, and the elements mapping of primary carbides in the Nb-1 sample is shown in Fig. 8. Elemental Nb and V are uniformly distributed in the primary carbide, while elemental Ti is mainly enriched in the central portion, and elemental Mo and Cr are more concentrated at the edges than at the core. Combined with the results in Fig. 5, it is well known that the Ti-rich carbide precipitates first during the solidification, and followed by V-rich and Nb-rich carbides. The rich-Mo-Cr carbide finally precipitates on the periphery of primary carbide. The Ti-rich carbide can act as the nucleation core of the V-rich and Nb-rich carbides. Considering that the Ti-rich carbide locates at the core, while their size is small, it is reasonable that the Ti-rich carbide has not been found in the 2D observation.

Fig. 8.

Elements mapping of primary carbide in Nb-1 sample. (Online version in color.)

The variation of Nb content and cooling rate with number density and size of primary carbides in the 3D observation are shown in Figs. 9 and 10. The error bars represent the standard deviation. As the Nb content increases, the number density of primary carbides increases from 71 to 83.75 mm⁻². However, the number density of primary carbides in the Nb-4 sample is only 21.1 mm⁻², which is much lower than that of the Nb-3 sample, indicating that the cooling rate also has a great influence on the number density of primary carbides. As the increase of Nb content, the average size of primary carbides increases rapidly from 31.71 to 42.93 μm. Besides, the primary carbide size in the Nb-4 sample reaches 86.32 μm, indicating that a faster cooling rate can effectively decrease the size of primary carbides. The addition of elemental Nb increases the number density and size of primary carbides, which in turn strengthens the detriment of primary carbides to H13 steel. Therefore, the appropriate Nb content in H13 steel should be chosen by synthetically considering the effect of Nb content on primary carbides and the material performance. Although the low cooling rate can effectively reduce the number density of primary carbides, it significantly increases the size of primary carbides. When the cooling rate is faster, the crystal grains have a higher nucleation rate and a faster growth rate, providing less space for the growth of primary carbides. However, when the cooling rate is slower, the crystal grains have a lower nucleation rate, which provides enough space for the growth of primary carbides. Therefore, the primary carbide size cooled in air is much smaller than cooled in furnace. The detriment of primary carbide size is larger than the number density, so the rapid cooling solidification conditions should be adopted in the actual smelting process.

Fig. 9.

Variation of Nb content with number density of primary carbide in 3D observation. (Online version in color.)

Fig. 10.

Variation of Nb content with primary carbide size in 3D observation. (Online version in color.)

4. Discussion

4.1. Formation Process of the Last-to-solidify Region

Since the microsegregation of alloying elements will occur during the solidification, the Scheil solidification model was used to calculate the solidification process of Nb-microalloying H13 steel. Taking the chemical composition of the Nb-2 sample in Table 1 as an example, the mass percentage of main elements in Liquid during the solidification is shown in Fig. 11. At the first stage of solidification, the content of Cr, Mo, V, C, and Nb remains unchanged. However, the alloying elements content increase rapidly at the end of solidification, especially the elemental Nb, Mo, and V. Therefore, the content of alloying elements along the grain boundaries is far larger than in the matrix, finally leading to the formation of network last-to-solidify region. When the enriched content of elemental Cr, Mo, V, C, and Nb reaches the thermodynamic precipitation conditions of the corresponding primary carbide, the primary carbide begins to precipitate and then rapidly grows into dendritic structure in last-to-solidify region. Therefore, the primary carbide and enrichment of alloying elements only can be found in the last-to-solidify region, as shown in Fig. 12, which agrees well with the calculated results. The segregation of elemental Cr, Mo, V, C, and Nb during solidification is the main reason for the precipitation of primary carbides in the last-to-solidify region. Moreover, the large primary carbide cannot be completely eliminated just by heat treatment because of its high-temperature thermal stability, further leading to the reduction of the material performances.²²⁾

Fig. 11.

Mass percent of main elements in Liquid during solidification. (Online version in color.)

Fig. 12.

Elements mapping of the last-to-solidify region and primary carbide analyzed by EPMA in Nb-2 sample. (Online version in color.)

4.2. Effect of Nb Content on the Precipitation Process of Primary Carbides

Figure 13 shows the precipitation process of the primary carbides during solidification in all samples, which was calculated with the Scheil model. The primary carbides and MnS precipitate at the end stage of solidification. Taking the Nb-2 sample as an example, it is well known that the Ti4C2S2 phase enriched in elemental Ti, C, and S precipitates first and followed by MnS inclusion. As the solid fraction increases, the precipitation sequence of carbides is Mo-rich, Cr-rich, and V-rich phases. Moreover, the Ti4C2S2, Mo-rich, Cr-rich, MnS, and V-rich phases exist until the liquid phase completely solidifies. The Nb content has a great influence on the precipitation sequence and the solid fraction where primary carbide starts to precipitate. As shown in Table 2, as the Nb content increases, the precipitation solid fraction of Ti4C2S2 phase is almost unchanged, and the precipitation solid fraction of Mo-rich carbide decreases rapidly even smaller than the Cr-rich carbide. The increase of Nb content promotes the precipitation of primary carbides during solidification, especially the Mo-rich carbide, which will provide better kinetic conditions for the nucleation and growth of the primary carbides.

Fig. 13.

Precipitation processes of precipitated phases calculated with Scheil model. (a) Nb-0 sample, (b) Nb-1 sample, (c) Nb-2 sample, (d) Nb-3 sample. (Online version in color.)

Table 2. Solid fraction at which primary carbide begins to precipitate.

Nb content	solid fraction
	Ti4C2S2	MnS	Cr-rich	Mo-rich	V-rich
0	0.8777	0.9172	0.9292	0.9419	0.9523
0.02%	0.8777	0.9128	0.9286	0.9418	0.9524
0.04%	0.8728	0.9190	0.9386	0.9255	0.9509
0.08%	0.8721	0.9168	0.9376	0.9244	0.9512

According to the experimental results in Fig. 8 and thermodynamic calculated results in Fig. 13, we can obtain the precipitation mechanism of primary carbides in Nb-microalloying H13 steel. The Ti4C2S2 phase precipitates first in the last-to-solidify region and then quickly grows into the dendritic structure. The Mo-rich and/or Cr-rich carbides (Mo-Cr-rich carbide) are formed together with the Ti4C2S2 phase as the nucleation core. The V-rich carbide just starts to precipitate when the solid fraction is 0.95, and the corresponding temperature at this time is only 1160°C, so the V-rich carbide may not be able to precipitate during the actual solidification of H13 ingots. In addition, what we found at the center of the primary carbides is the enrichment of elemental Ti, Nb, and V instead of the Ti4C2S2 phase. Therefore, the equilibrium model was used to calculate the transformation process of Ti4C2S2 phase during the cooling process, and the results are shown in Fig. 14. Taking the Nb-2 sample as an example, as the temperature decreases, the Ti4C2S2 phase first transforms into Nb-rich carbide. Since the temperature at this time is higher and the diffusion rate of elements in the solid phase is larger, the solid phase transformation is relatively sufficient, finally resulting that the untransformed Ti-rich carbide only can be found in the core of primary carbides. As the temperature further decreases, the Nb-rich carbide partly transforms into V-rich carbide because of the lower diffusion rate of elements in the solid phase. Therefore, the distribution of elemental V and Nb in the primary carbide is the same. When the temperature decreases to the room temperature, the elements mapping results in Fig. 8 was formed. According to Xie’s research results,²³⁾ the Ti-V-rich phase precipitates first during solidification of Nb microalloying H13 steel, after which Nb-rich phase and V-rich precipitates successively, which is somewhat different from the precipitation mechanism in this article. The difference in S content may be the main reason for the different precipitation mechanisms of the primary carbides.

Fig. 14.

Transition processes of Ti4C2S2 phase calculated with equilibrium model. (a) Nb-0 sample, (b) Nb-1 sample, (c) Nb-2 sample, (d) Nb-3 sample. (Online version in color.)

The Nb content has a great influence on the transformation process of Ti4C2S2 phase. As the increase of Nb content, the transformation temperature of Ti4C2S2 phase to Ti(C, N) or Nb-rich carbide increases rapidly, which provides better kinetic conditions for the formation of Nb-rich carbide. Therefore, the proportion of the mass fraction of Nb element in the Nb-containing primary carbide gradually increases, which is completely consistent with the experimental results in Fig. 6. According to Guo’s research results,²⁴⁾ the smaller the cooling rate, the larger the primary carbide size. Therefore, the primary carbide size in the Nb-4 sample is much larger than in the Nb-3 sample.

4.3. Effect of Nb Content on the Growth Behavior of Primary Carbides

Since the growth of primary carbides is largely depends on the kinetic conditions, the effect of Nb content on the growth behavior of the primary carbides was discussed in this part. Ohnaka’s equations are used to calculate the segregation of elements during the solidification.²⁵⁾   

$$C_L^i=C_0^i\cdot {\left(1-\left(1-\beta \cdot \frac{k_i}{\left(1+\beta \right)}\right)\cdot f_s\right)}^{\frac{\left(k_i-1\right)}{\left(1-\beta \cdot \frac{k_i}{\left(1+\beta \right)}\right)}}$$(1)

  

$$\beta =4\cdot D_S^i\cdot \frac{t_f}{{\lambda }_2{}^2}$$(2)

Where, $C_L^i$ is the content of the solute i in the liquid, $C_0^i$ is the initial content of the solute i, k_i is the partition coeffient of solute i, f_s is the solid fraction. $D_S^i$ is the solute diffusivity of solute i in the solid phase, m²/s, λ₂ is the secondary dendrite arm spacing, m; t_f is the local solidification time, s, which can be obtained by Eq. (3).   

$$t_f=\frac{T_L-T_S}{R_C}$$(3)

Where T_L and T_S are the liquidus and solidus temperatures, K, obtained from the calculation results of Thermodynamic software. R_C is the local cooling rate, K/s, which can be calculated by Eq. (4).²⁴⁾   

$${\lambda }_2=175.4\cdot R_C{}^{-0.322}$$(4)

The temperature of the solidification interface, T, can be calculated as follows:   

$$\text{T}={\text{T}}_0-\frac{T_0-T_L}{1-f_S\cdot \frac{T_L-T_S}{T_0-T_S}}$$(5)

Since the experimental process of the Nb-0, Nb-1, Nb-2, Nb-3 samples is exactly the same, the cooling rates (R_C) are the same. As can be seen from Eq. (4), the secondary dendritic spacing (λ₂) of the Nb-0, Nb-1, Nb-2, and Nb-3 samples is the same. Moreover, considering that the main components of the Nb-0, Nb-1, Nb-2, and Nb-3 samples are also basically the same, the solid and liquid temperatures (T_L and T_S) of the four samples are the same. Therefore, the segregation content of each element ($C_L^i$) in Nb-0, Nb-1, Nb-2, Nb-3 samples is the same.

During the solidification process, the Ti4C2S2 phase precipitates first, and then Mo-Cr-rich carbide precipitates. Therefore, the following equilibriums were assumed to exist on the interface of the precipitated primary carbide and the molten steel.   

$$\begin{array}{l}4\left[\mathrm{T}\mathrm{i}\right]+2\left[\mathrm{C}\right]+2\left[\mathrm{S}\right]=\mathrm{T}\mathrm{i}4\mathrm{C}2\mathrm{S}{2}_{(\mathrm{s})}\\ 2\left[\mathrm{M}\mathrm{o}\right]+\left[\mathrm{C}\right]=\mathrm{M}\mathrm{o}2{\mathrm{C}}_{(\mathrm{s})}\end{array}$$

Since the primary carbides precipitate at the end of solidification, the interfacial reaction is comparatively fast. In addition, the Ti content is relatively lower than elemental C and S, and the C content is lower than Mo. Therefore, the mass transport of Ti for the Ti4C2S2 phase and C for the Mo2C carbide is assumed as the controlling step during the growth of the primary carbides. The growth of the primary carbides is expressed by Eqs. (6) and (7).   

$$\begin{array}{l}{\text{r}}_{Ti4C2S2}=\\ {\left({\displaystyle \frac{M_{Ti4C2S2}\cdot {\rho }_{Fe}}{50\cdot M_{Fe}\cdot {\rho }_{Ti4C2S2}}}\cdot D_L^{Ti}\cdot \left(C_L^{Ti}-C_{eq}^{Ti}\right)\cdot \left(1-f_{s,Ti4C2S2}\right)\cdot t_f\right)}^{0.5}\end{array}$$(6)

  

$${\text{r}}_{Mo2C}={\left(\frac{M_{Mo2C}\cdot {\rho }_{Fe}}{50\cdot M_{Fe}\cdot {\rho }_{Mo2C}}\cdot D_L^C\cdot \left(C_L^C-C_{eq}^C\right)\cdot \left(1-f_{s,Mo2C}\right)\cdot t_f\right)}^{0.5}$$(7)

Where, r is the primary carbide radius, m; ${\text{D}}_L^{Ti\mathrm{\hspace{0.17em} } or\mathrm{\hspace{0.17em} } C}$ is the diffusion coefficient of solute Ti or C in molten steel, m²/s; M is the molecular weight; ρ is the density, kg/m³; ${\text{C}}_{eq}^{Ti\mathrm{\hspace{0.17em} } or\mathrm{\hspace{0.17em} } C}$ is the equilibrium content of Ti or C elements. f_{s,Ti4C2S2 or Mo2C} is the solid fraction that Ti4C2S2 or Mo2C precipitate.

According to above calculated results, the $C_L^{Ti\mathrm{\hspace{0.17em} } or\mathrm{\hspace{0.17em} } C}$ in Nb-0, Nb-1, Nb-2, and Nb-3 samples are the same, so the corresponding ${\text{C}}_{eq}^{Ti\mathrm{\hspace{0.17em} } or\mathrm{\hspace{0.17em} } C}$ are also the same. Therefore, it can be seen from Eqs. (6) and (7) that the size of r is only positively related to the value of (1−f_{s,Ti4C2S2 or Mo2C}). According to the calculated results in Table 2, the size of Ti4C2S2 phase in Nb-0, Nb-1, Nb-2, Nb-3 samples is basically the same and the size of Mo-Cr-rich carbide gradually increases with the increase of Nb content. Therefore, the size of the final primary carbides increases with the increase of Nb content, which is consistent with the statistical results in Fig. 10.

5. Conclusions

(1) There is a huge difference between the 2D and 3D morphologies of the primary carbide. The enrichment of the elemental Cr, Mo, V, and C in the last-to-solidify region leads to the formation of the primary carbides. As the Nb content increases, the last-to-solidify region size gradually decreases.

(2) The increase of Nb content has little influence on the 2D and 3D morphologies of the primary carbide. However, as the Nb content increases, the number density and size of primary carbide in the 3D observation gradually increase. The lower cooling rate reduces the number density of the primary carbide but significantly increases the size.

(3) During the solidification of Nb microalloying H13 steel, the Ti4C2S2 phase precipitates first, and then the Mo-rich or Cr-rich carbides precipitate outside the Ti4C2S2 phase. Moreover, during the cooling process of H13 steel, the Ti4C2S2 phase transforms into Nb-rich and V-rich carbides. The increase of Nb content promotes the precipitation of primary carbides during solidification and phase transition of Ti4C2S2 phase after solidification, resulting in an increase in the number density and size of the primary carbides.

Acknowledgements

The authors are grateful for support from the National Natural Science Foundation of China (NO. 51874034).

References

• 1)   G. H.  Yan,  X. M.  Huang,  Y. Q.  Wang,  X. G.  Qin,  M.  Yang,  Z. M.  Chu and  K.  Jin: Met. Sci. Heat Treat., 52 (2010), 393.
• 2)   J. H.  Lee,  J. H.  Jang,  B. D.  Joo,  Y. M.  Son and  Y. H.  Moon: Trans. Nonferrous Met. Soc. China, 19 (2009), 917.
• 3)   M.  Pérez and  F. J.  Belzunce: Mater. Sci. Eng. A, 624 (2015), 32.
• 4)   J. R. C.  Guimarães,  J. R. T.  Branco and  T.  Kajita: Mater. Sci. Technol., 2 (1986), 1074.
• 5)   W. B.  Morrison and  J. H.  Woodhead: J. Iron Steel Inst., 201 (1963), 43.
• 6)   W. B.  Morrison: J. Iron Steel Inst., 201 (1963), 317.
• 7)   J. R. T.  Branco and  G.  Krauss: Heat Treatment and Surface Engineering: New Technology and Practical Applications, ed. by G. Krauss, ASM International, Materials Park, (1988), 195.
• 8)   C. N.  Elias and  C. S.  Da Costa Viana: J. Mater. Eng. Perform., 1 (1992), 751.
• 9)   C. N.  Elias and  C. S.  Da Costa Viana: Mater. Sci. Technol., 8 (1992), 785.
• 10)   J. L.  Maloney,  W. P.  Edwards and  M. S.  Rodney: Chromium Hot Work Steel, U.S. Patent No. 5207843, (1993).
• 11)   Y.  Xie,  G. G.  Cheng,  L.  Chen,  Y. D.  Zhang and  Q. Z.  Yan: Int. J. Miner. Metall. Mater., 23 (2016), 1264.
• 12)   Y.  Xie,  G. G.  Cheng,  L.  Chen,  Y. D.  Zhang and  Q. Z.  Yan: ISIJ Int., 56 (2016), 995.
• 13)   Y.  Xie,  G. G.  Cheng,  L.  Chen,  Y. D.  Zhang and  Q. Z.  Yan: Metall. Res. Technol., 113 (2016), 206.
• 14)   Q. C.  Jiang,  H. L.  Sui and  Q. F.  Guan: ISIJ Int., 44 (2004), 1103.
• 15)   J. T.  Tchuindjang and  J.  Lecomte-Beckers: Int. J. Fatigue, 29 (2007), 713.
• 16)   Y.  Huang,  G. G.  Cheng,  S. J.  Li and  W. X.  Dai: Steel Res. Int., 90 (2019), 1900035.
• 17)   J.  Yu and  C. J.  McMahon: Metall. Trans. A, 11 (1980), 277.
• 18)   P.  Doig and  P. E. J.  Flewitt: Acta Metall., 29 (1981), 1831.
• 19)   M. K.  Miller,  P.  Pareige and  M. G.  Burke: Mater. Charact., 44 (2000), 235.
• 20)   S. H.  Song,  R. G.  Faulkner and  P. E. J.  Flewitt: Mater. Sci. Eng. A, 281 (2000), 23.
• 21)   N.  D’Souza,  M.  Lekstrom and  H. B.  Dong: Mater. Sci. Eng. A, 490 (2008), 258.
• 22)   Y.  Xie,  G. G.  Cheng,  X. L.  Meng and  Y.  Huang: Metall. Res. Technol., 114 (2017), 206.
• 23)   Y.  Xie,  G. G.  Cheng,  L.  Chen,  Y. D.  Zhang and  Q. Z.  Yan: Steel Res. Int., 88 (2017), 1600119.
• 24)   M. T.  Mao,  H. J.  Guo,  F.  Wang and  X. L.  Sun: ISIJ Int., 59 (2019), 848.
• 25)   I.  Ohnaka: Trans. Iron Steel Inst. Jpn., 26 (1986), 1045.