The characteristics of precipitates larger than 1 µm in Nb–Ti microalloyed H13 tool steel were studied. Four types of large phases exist in the as-cast ingot according to the compositional characteristics, that is, (Ti,Nb,V)(C,N), V-rich carbide, Mo–Cr-rich carbide and sulfide. (Ti,Nb,V)(C,N) could be further classified as the Ti-rich one, the Nb-rich one and the Ti/Nb transforming one. V-rich carbide normally has a strip shape, while Mo–Cr-rich carbide presents a eutectic appearance. The compositional characteristics of large precipitates have little relation with the sample positions. V-rich and Mo–Cr-rich carbide have the tendency to dissolve at 1000°C and 1100°C, while the (Ti,Nb,V)(C,N) is more stable and there are no apparent changes on the morphology after holding 6 h at 1000°C and 3 h at 1100°C. These precipitates are generated during solidification. The precipitating process of (Ti,Nb,V)(C,N) could be well speculated through Thermo-Calc software. Ti-rich carbonitride precipitates first. Nb-rich carbonitride appears later, singly or on the Ti-rich carbonitride. V-rich carbide precipitates at the end of solidification. Eutectics consisting of iron matrix and Mo–Cr-rich carbide will be generated when the solutes in liquid steel reach the eutectic point.

AISI H13, one common tool steel with an excellent performance of hot strength and toughness, is widely used for hot extrusion, forging and die casting. Up to now, H13 is further improved, not only on the process optimization including the ESR process, but also in the field of composition adjustment according to the serving conditions.^1,2,3,4,5)

Several researchers have studied the effects of partial or total substitution of niobium for vanadium in H13 tool steel. In 1988, J. R. T. Branco et al.⁶⁾ replaced all of V by 0.07%Nb in H13, and found that the grain growth and the coarsening behavior of MC carbide was inhibited effectively. C. N. Elias and C. S. Da Costa Viana⁷⁾ reported their studies about the effect of partial substitution of 0.07%Nb for V on austenite grain size and hardness of H13 steel, and the results showed that the H13+Nb steel had a smaller austenite grain size and a finer carbide size distribution than the vanadium steel, which indicated a higher potential toughness. Shahram Kheirandish et al.⁸⁾ studied the effect of niobium addition on the microstructure of cast AISI hot work tool steel. The result showed that the 1.1% niobium addition by replacing 0.5% V and 0.4% Mo modified the cast structure of Nb-alloyed hot work tool steel, reduced the size and volume of eutectic cells, and increased the maximum hardness of the steel. James L. Maloney et al. applied for a patent⁹⁾ for a modified H13 hot working die steel named XP-2599 with 0.035% Nb and 0.38%V, that is, a partial substitution of niobium for vanadium. The modified steel showed an improved high temperature impact toughness and thermal fatigue resistance. In addition, they found that the addition of 0.01–0.2%Ti would further improve the property of steel product.

However, large precipitates could be found in the Nb microalloyed^10,11,12) or Nb+Ti^{13,14,15,16,17)} microalloyed steel and are stable even at high temperature, affecting the properties of steel products by wrecking the substrate continuity. These precipitates already exist in the as-cast ingot and normally precipitate during solidification. However, the generating mechanism of the large phases in microalloyed steel is still unclear even though some analyses have been taken in the previous study.^{13,14,15,16,17)} The paper is concerned on the characteristics and generating mechanism of large phases in the as-cast ingot of Nb–Ti microalloyed H13 tool steel, laying foundations for the further control of the large phases.

One ingot of approximately 5 kg weight with circular truncated cone shape (the bottom radius 60 mm, the top radius 80 mm, the height 145 mm) was manufactured in vacuum induction furnace according to the chemical composition requirements in JIS G4404-2006. The iron was first melted, and then the graphite block and several alloys with high purity were added to alloy the molten steel. After the homogenization of composition, the molten steel was casted into a cast iron mould with the size φ145 mm×170 mm. The ingot was then stripped and air cooled to the room temperature.

The average ingot composition is listed in Table 1. 0.036 wt% Ti and 0.064 wt% Nb are contained in the ingot. The content of O and N is 23 ppm and 34 ppm respectively. In addition, 48 ppm Al and 46 ppm S from the impurities in alloys exist in the ingot.

The longitudinal section of the ingot through diameter is shown in Fig. 1. One shrinkage cavity exists at the top. One sample with height 53 mm was cut to observe the solidification microstructure after etching by the hydrochloric acid water solution with the volume ratio 1:1 at 70–80°C. The dotted box region with the size 53 mm×46 mm was detected by original position analyzer for metal (OPA-200) to analyze the macrosegregation of alloying elements. Three cubic samples (numbered as 1, 2 and 3) with the side length 15 mm were cut in 1/2 radius and center on the cross section in 1/2 height. SEM equipped with EDS was applied to observe the characteristics of the large precipitates in the samples polished by SiC papers and polishing paste. The precipitates larger than 1 μm were recorded. In addition, the samples were etched by 4% nital to determine the sites of large precipitates.

Fig. 1.

The sampling method.

The solidification microstructure of the ingot is captured by a scanner and shown in Fig. 2. An obvious porosity region exists in the center. The microstructure in position 1, 2 and 3 is marked by the box. Fine columnar crystals develop towards the center in position 1 and 3, while fine equiaxed grains exist in position 2. In addition, ‘V’ type segregation can be observed at the position 2.

Fig. 2.

The solidification microstructure of ingot.

The cooling rate during solidification is predicted by Eq. (1).¹⁸⁾ λ_SDAS is the secondary dendrite arm spacing. C_R is the cooling rate (°C/s) and C_C is the carbon content (weight pct). The average secondary dendrite arm spacing for position 1, 2 and 3 is 41.36 μm, 45.25 μm and 41.87 μm respectively after etching by picric acid at 60–70°C and observed by optical microscope, as shown in Fig. 3. The calculated value for cooling rate in the position 1, 2 and 3 is 48.4°C/s, 37.7°C/s and 46.8°C/s respectively.   

Fig. 3.

The secondary dendrite arm spacing for the three samples, (a) position 1, (b) position 2, (3) position 3.

The macrosegregation of the matrix elements, C, Si, Mn, Cr, Mo and V on the selected section is shown in Fig. 4. The numbers represent the weight percent content of the corresponding elements. There exists an obvious rich region for C in the central area. The maximum content is 0.53 wt%. For Si, Mn, Cr, Mo and V, the compositional distribution is relatively uniform.

Fig. 4.

The macrosegregation in the longitudinal section.

The large precipitates observed in the samples can be classified as follows according to the compositional characteristics.

(1) (Ti,Nb,V)(C,N). These precipitates mainly contain the alloying elements Ti, Nb and V. The normalized compositional distribution of Ti, Nb and V is shown in Fig. 5. According to the compositional characteristic, it could be further classified into the Ti-rich one, the Nb-rich one and the Ti/Nb transforming one which is marked as ‘Ti’, ‘Nb’ and ‘Ti-Nb’ respectively. The compositional characteristics in position 1, 2 and 3 are consistent with each other. It shows that the cooling rate and solidification structure in the different position of the ingot in the present situation has little effect on the compositions of these precipitates. It may be due to the low diffusion amount of the carbonitride forming elements in solid steel under the high cooling rate.

Fig. 5.

The normalized atomic ratio of Ti, Nb and V in (Ti,Nb,V)(C,N) type phase.

The typical morphology of ‘Ti’ and ‘Nb’ is shown in Figs. 6(a) and 6(b) respectively. The corresponding atomic ratios of alloying elements analyzed by EDS are listed in Table 2. Ti-rich phase is a carbonitride and has a polygon shape. The size is mostly smaller than 5 μm. The contents of Cr and Mo are very low. Nb-rich phase is relatively larger and mainly in the form of carbide. A certain ratios of Ti and V are contained and the contents of Mo and Cr are also negligible. In addition, some large phases precipitate on the black oxides the composition of which are mainly (Al,Ti)O. Heterogeneous nucleation could reduce the nucleation barrier and contribute the precipitation.^19,20,21)

Fig. 6.

The typical morphology of (a) Ti-rich and (b) Nb-rich (Ti,Nb,V)(C,N) type phase.

For Ti/Nb transforming phase which is shown in Fig. 7, the size is even up to 10 μm. The line scanning mapping of one typical precipitate is shown in Fig. 8. With the region shifting from the dark gray one to the white one, the content of Ti decreases gradually and that of Nb increases, while the content of V has little change.

Fig. 7.

The typical morphologies of Ti/Nb transforming phase.

Fig. 8.

Line scanning mapping of Ti/Nb transforming phase.

(2) V-rich carbide. The typical morphologies of V-rich carbides in three samples are shown in Fig. 9 which is marked as ‘V’. V-rich carbide normally has a strip shape the size of which is mostly larger than 5 μm. The compositions in the sites the arrows pointing to are listed in Table 3. A certain content of Mo, Nb and Cr are contained, while the content of Ti is negligible. There exist little difference for the morphology and composition of V-rich carbide in the different sample positions.

Fig. 9.

The typical morphology of V-rich and Mo–Cr-rich carbides in (a) position 1, (b) position 2 and (c) position 3.

(3) Mo–Cr-rich carbide. As Fig. 9 shown with the symbol ‘Mo–Cr’, Mo–Cr-rich carbide has a eutectic appearance the size of which could be up to tens of microns. The compositions are listed in Table 3. The main elements are Mo and Cr. The content of both has little difference. A relatively lower content of V are contained and the contents of Nb and Ti are negligible. In addition, Nb-rich carbide, V-rich carbide and Mo–Cr-rich carbide always connect with each other end to end, which may be related with the precipitation sequence. Similar to other large phases described above, the difference for the characteristic of Mo–Cr-rich carbides in different positions is not obvious.

(4) Sulfide. As shown in Figs. 6(b) and 9(b) with the symbol ‘S’, sulfide normally has an elliptical shape. Most of the sulfide is smaller than 5 μm.

The etching pattern of the sample observed by SEM is shown in Fig. 10(a). Figure 10(b) is the partial enlargement of the box region. The gray area is the boundary of the dendritic crystal in which the alloying elements are richer than the matrix due to the segregation during solidification. The black line is the grain boundary. The large precipitates are marked according to the compositional characteristics. It shows that the large precipitates are located in the boundary of the dendritic structure. In other words, these precipitates precipitate in the liquid steel during solidification. Moreover, Mo–Cr-rich carbides are mostly located in the joint area of the multiple crystals which is the final solidification region. It may precipitate finally compared with other types of precipitates.

Fig. 10.

(a) The precipitating site of the large precipitates, (b) the partial enlargement for the box region.

The thermal stability of the large precipitates is assessed by quenching from 1000°C after holding for 0.5 h, 3 h, 6 h and 1100°C for 0.5 h and 3 h. The results are shown in Figs. 11 and 12. The temperature and time are located on the top right corner.

Fig. 11.

The morphologies of large phases after holding 0.5 h, 3 h and 6 h at 1000°C.

Fig. 12.

The morphologies of large phases after holding 0.5 h and 3 h at 1100°C.

V-rich and Mo-Cr-rich carbides are unstable at 1000°C and 1100°C. Mo-Cr-rich carbide could not be observed in the samples after 0.5 h at 1100°C. V-rich carbide with strip shape also has the tendency to dissolve and starts to fracture after 0.5 h, as shown in Figs. 11(a) and 12(a). After 6 h, there are no large V-rich and Mo–Cr-rich carbides left in the sample. However, for (Ti,Nb,V)(C,N), the change of the morphologies could not be observed. The corresponding compositional distributions after different times are shown in Figs. 13(a), 13(b) and 13(c) respectively. The change of composition is also not obvious even though the large precipitates observed in the samples are not equilibrium ones in the solid steel after solidification. Actually, the change of composition could not be observed for that the composition tested is just an average one and the changes for the precipitates with large size may be minor under the present testing method.

Fig. 13.

The compositional distribution of (Ti,Nb,V)(C,N) after holding different times at 1000°C and 1100°C, (a) Ti-rich phase, (b) Ti/Nb transforming phase, (c) Nb-rich phase.

During solidification, microsegregation will occur. In position 1 and 3, with the development of columnar crystal, the solute is gradually rich in liquid steel between the secondary dendrite arms, resulting in the precipitation of primary phases when the contents of alloying elements reach the precipitating condition.

The microsegregation and precipitating process of the primary phases in position 1 and 3 can be speculated through the Scheil-Gulliver model in Thermo-Calc software for the limited diffusion amount of the carbonitride forming elements in solid steel under the high cooling rate in the present situation. The model assumes that the alloying elements are homogeneous in liquid steel, the diffusion is not occurred in solid steel and local equilibrium at solid-liquid interface is maintained during solidification.

In Scheil-Gulliver solidification, the new composition of the liquid can be determined by making a stepping operation on the temperature variable with small decrementing steps. After each step, the amount of formed solid phase is removed and the overall composition is reset to the new liquid composition. During calculation, the final composition of liquid steel before solidification was first determined and listed in Table 4. The content of of O, Al and Ti left in liquid steel is 0.00032 wt%, 0.0029 wt% and 0.0355 wt% respectively for the generation of oxide (Al,Ti)O. The content of O is very low. During calculation of solidification process, the composition in Table 4 was used as the initial composition of liquid steel. The content of O was ignored, or some unreasonable oxide would appear in Scheil Model.

The calculation result is shown in Fig. 14. Seven phases will appear during solidification, that is, the matrix ferrite phase BCC_A2#1, the matrix austenite phase FCC_A1#1, and the primary precipitates FCC_A1#2, M₇C₃, M₆C, MnS and Ti₄C₂S₂, apart from the oxide (Al,Ti)O which already exists in the liquid steel before solidification. During solidification, the primary precipitate FCC_A1#2 first appears and then Ti₄C₂S₂, M₆C and M₇C₃ precipitates successively. MnS does not precipitate until solidification process reaches to the end. The corresponding solid fractions of precipitation for the primary phases are listed in Table 5. Compared with the connecting forms of large precipitates in Fig. 9, the precipitation sequence is consistent with each other.

Fig. 14.

Temperature as a function of solid fraction during solidification.

The mole fraction of alloying elements in FCC_A1#2, M₆C and M₇C₃ is shown in Figs. 15 and 16 respectively. There is no value at fraction of solid 1.0 for that the final solute concentration in liquid steel will become infinite and the value is insignificant in Scheil model. According to the compositional characteristics of FCC_A1#2 phase, Ti-rich carbonitride precipitates firstly when the solid fraction reaches 0.704. Subsequently, Nb-rich carbide with a certain content of Ti and V will precipitate. V-rich carbide appears finally. A certain content of Mo, Cr and Fe are contained. The content of Nb is low. M₆C is a Mo–Fe-rich carbide phase with a certain content of Si and little content of V and Cr, while M₇C₃ is a Cr–Fe-rich carbide phase and has a little content of V and Mo.

Fig. 15.

Mole fraction of alloying elements as a function of solid fraction, (a) FCC_A1#2, (b) the sectional representation.

Fig. 16.

Mole fraction of alloying elements as a function of solid fraction, (a) M₆C phase, (b) M₇C₃ phase.

The compositional evolution of FCC_A1#2 seems to be consistent with that of (Ti,Nb,V)(C,N) observed in the samples. To compare with the experimental results in Fig. 5, the normalized atomic ratio of Ti, Nb and V in FCC_A1#2 is plotted in the Ti–Nb–V ternary phase diagram, as shown in Fig. 17. The empty triangle represents the calculated results and the curve with an arrow represents the compositional changing direction with the development of solidification. The region without date could be speculated according to the changing tendency of elements Ti, Nb and V and is presented as the dotted line. The compositional distribution of (Ti,Nb,V)(C,N) type phase in Fig. 5 is also added into the diagram and presented as the half-solid circle. The calculated results agree with the measured data well. The generating mechanism of (Ti,Nb,V)(C,N) type phase can be well speculated by the behavior of FCC_A1#2 phase.

Fig. 17.

The compositional comparison between calculated results and experimental results for (Ti,Nb,V)(C,N).

According to the calculation results described above, V-rich carbide FCC_A1#2, M₆C, M₇C₃ and sulfide will precipitate successively at the end of solidification. The compositional characteristics of V-rich region in FCC_A1#2 are consistent with that in the V-rich carbide observed in the samples. For Mo–Cr–rich carbide, it might be the mixture of M₆C and M₇C₃. Actually, the generating mechanism of Mo–Cr-rich carbide in Fe–V–Cr–Mo–C system is more complex.^22,23) When the enrichments of solutes in the liquid steel reach the eutectic point, Mo–Cr-rich carbide will appear in the form of eutectic consisting of iron matrix and mixed (Mo,Cr,V,Fe)_xC_y.

Based on the analysis above, the generating mechanism of the large precipitates in position 1 and 3 could be described as follows. During solidification, the columnar structure gradually develops and the solutes C, N, Nb, Ti, V, Mo and Cr are rejected into the liquid steel. As shown in Fig. 18, Ti-rich (Ti,Nb,V)(C,N) precipitates first for the higher stability of TiN than other nitride and carbide.²⁴⁾ With the development of solidification, the contents of Nb and C in the liquid steel gradually increase. When the solid fraction reaches a certain value, Nb-rich (Ti,Nb,V)(C,N) starts to precipitate instead of Ti-rich one, singly or on the Ti-rich phase which has already existed in the liquid steel. V-rich carbide will precipitate at the end of solidification. With the enrichment of solutes in liquid steel, eutectic consisting of iron matrix and Mo–Cr-rich carbide will appear when the contents of alloying elements reach the eutectic point.

Fig. 18.

The generating process of large phase during solidification.

The segregation way in position 2 is more complex. V-segregates arise due to the fissuring of networks of loosely-connected equiaxed grains under action of metallostatic head and solidification shrinkage, which leads to the formation of open shear planes that can fill with any remaining liquid.^25,26) The remaining liquid will be enriched gradually with the development of solidification. According to the analysis for the composition distribution of large precipitates, the characteristics in position 2 are similar to that in position 1 and 3. It reveals that the enrichments of solutes can still reach the concentrations for precipitation of primary phases with the development of equiaxed crystals under the high cooling rate.

The large precipitates existing in the ingot certainly affect the quality of steel product. V-rich and Mo–Cr-rich carbides could be controlled easily by heat treatment at high temperature, while it will be useless for (Ti,Nb,V)(C,N). Taking reasonable measures to refine as-cast structure^27,28,29,30) may be more effective.

(1) Four types of large precipitates exist in the as-cast ingot, that is, (Ti,Nb,V)(C,N), V-rich carbide, Mo–Cr-rich carbide and sulfide. (Ti,Nb,V)(C,N) could be further classified as the Ti-rich one, the Nb-rich one and the Ti/Nb transforming one. V-rich carbide normally has a strip shape, while Mo–Cr-rich carbide presents a eutectic morphology. The size of both can be up to tens of microns. The compositional characteristics of large precipitates have little relation with the sample position.

(2) V-rich and Mo–Cr-rich carbide have the tendency to dissolve at 1000°C and 1100°C, while (Ti,Nb,V)(C,N) is more stable and there are no apparent changes on the morphology and even composition after holding 6 h at 1000°C and 3 h at 1100°C.

(3) The large precipitates precipitate during solidification. The generating process could be well speculated through Thermo-Calc. Ti-rich (Ti,Nb,V)(C,N) precipitates first. Nb-rich carbonitride will precipitate later, singly or on the Ti-rich carbonitride. V-rich carbide appears at the end of solidification. When the solute in liquid steel reaches the eutectic point, eutectics consisting of iron matrix and Mo–Cr-rich carbide will be generated.

The authors would like to express their sincere thanks to the Xining Special Steel Co. Ltd for the technical help.

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