Precipitation Behavior of Primary Precipitates in Ti-microalloyed H13 Tool Steel

Abstract

The characteristics and thermal stability of the primary precipitates in H13 ingot modified with 0.034 wt% Ti were studied. A large number of (Ti,V) carbonitrides were observed, including the homogeneous Ti-rich ones and the inhomogeneous Ti-V-rich ones. The atomic ratios of Ti/V in Ti-rich precipitates range from 3.36 to 9.53, while those in Ti-V-rich precipitates range from 0 to 7.60. A certain number of sulfide and Mo-Cr-rich carbides also exist in the ingot. Ti-rich carbonitride is stable at 1150°C and 1250°C. Ti-V-rich precipitates with low content of Ti are unstable and start to decompose after holding a short time at high temperature. No Mo-Cr-rich carbide is left after 3 h at 1150°C or even 0.5 h at 1250°C. These precipitates are generated during solidification and the generating process could be well speculated by Thermo-Calc. During solidification, Ti-N-rich carbonitride precipitates first and the composition changes little until the generation of Ti-V-rich precipitates. With the development of solidification, the contents of Ti and N in newly generated Ti-V-rich precipitates decrease, while that of V and C increase. Sulfide and Mo-Cr-rich carbides precipitate at the end of solidification.

1. Introduction

Titanium has been widely used as a microalloying element in steels to prevent the coarsening of austenite through a dispersion of fine TiN particles, such as the improvement of toughness of weld heat affected zone,^1,2,3,4,5) or increasing the fraction of equiaxed grains in solidification structure through the heterogeneous nucleation of δ-phase on TiN or TiN+oxide.^6,7,8) The selection and control of particle size is important to the utilization efficiency of titanium. However, large precipitates can be always found in the Ti-containing steels.^9,10,11) The size could even reach tens of microns. These precipitates will be detrimental to the qualities of steel products instead of playing any useful role.

Practical steels usually contain more than one microalloying elements besides titanium, and large complex precipitates instead of the pure ones will exist in the steels. Several studies have paid attention to the characteristics and thermal stability of the large precipitates in these steels. D.C. Houghton⁹⁾ studied the characteristics of carbonitride in HSLA steels with 0.011–0.039 wt% Ti and 0.050–0.054 wt% Nb. Dendritic niobium-rich and cuboid titanium-rich (Ti,Nb)N were observed, and the length size of titanium-rich one could be several microns. Niobium-rich (Ti,Nb)N were metastable at high temperature in austenite and would transform to an array of titanium rich cuboids on prolonged exposure at temperature above 1100°C. Z. Chen et al.¹⁰⁾ studied the characteristics of large precipitates in Nb–Ti–V, Nb–Ti and Nb–V HSLA steel, and also observed large cuboidal and dendritic precipitates in Nb–Ti–V steel ingot, while only dendritic precipitates existed in Nb–Ti steels. The cuboidal precipitates were rich in Ti and N, and were stable up to the melting point of the steel. The dendritic precipitates had various contents of Nb and Ti. The ones with high content of Nb were less stable and would break up into many small parts at high temperature. No large precipitates were found in the Nb–V steel. It was thought that the addition of titanium to niobium-containing HSLA steels resulted in undesirable precipitation. M. Prikryl¹¹⁾ et al. observed large dendritic V-rich (Ti,V)N precipitates with a wide range of Ti contents in a medium carbon and Ti–V–N microalloyed steel ingot. These precipitates would transform to equiaxed, cuboidal precipitates with higher content of Ti after heat treatment. The large precipitates all above are normally considered to be generated in the liquid steel due to the segregation of solute during solidification. However, the detailed generating mechanism is still unclear.

AISI H13 is a medium carbon steel with about 0.4 wt% C, widely used for hot extrusion, forging and die casting.¹²⁾ About 8 wt% carbide forming elements, such as V, Mo and Cr, are contained to prevent the coarsening of austenite or strength the steel through secondary hardening. However, large precipitates with micron size are often observed such as (Ti,V) carbonitride in as-cast ingot or even bar after heating treatment.^13,14,15) The element titanium is most from the impurity in alloys during the production process. In addition, several studies on the adjustment of composition of H13^16,17,18) pointed out that the addition of a certain content of titanium was helpful to further improve the property of steel product. This paper is concerned on the characteristics and further thermal stability of large precipitates in the as-cast ingot of H13 tool steel modified with 0.034 wt% titanium, laying foundations for the control of these precipitates.

2. Experimental Method

One ingot of approximately 5 kg weight was manufactured in a vacuum induction furnace according to the chemical composition requirements in JIS G4404-2006. Iron was first melted, and then alloyed by graphite block and several alloys with high purity. The raw material for titanium was titanium sponge. After homogenization of the 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 composition of the ingot is listed in Table 1. 0.034 wt% Ti is contained. The content of O and N is 0.0019 wt% and 0.0033 wt% respectively. The total content of aluminum Alt is 0.0074 wt% and the soluble one Als is 0.0053 wt%. In addition, 0.0059 wt% S from the impurities in alloys exists in the ingot.

Table 1. Composition of ingot, wt%.

C	Si	Mn	Cr	Mo	V	Ti	Als	Alt	S	O	N
0.35	1.03	0.40	5.29	1.51	0.97	0.034	0.0053	0.0074	0.0059	0.0019	0.0033

The morphology of the ingot after cut through diameter is shown in Fig. 1. The ingot has a circular truncated cone shape with bottom diameter 60 mm, top diameter 79 mm and height 145 mm. An obvious shrinkage cavity exists at the top. One sample with height 67 mm was cut to observe the solidification microstructure after etching by the hydrochloric acid water solution with volume ratio 1:1 at 70–80°C. Three cubic samples with side length 15 mm were cut in the edge and center on the cross section and numbered as 1, 2 and 3. A FEI Quanta scanning electron microscope with an EDS x-ray analyzer was applied to observe the characteristics of the precipitates larger than 1 μm in the samples after grounded by SiC papers and polished by polishing paste. In addition, the samples were further etched by 4% nital at room temperature to determine the secondary dendrite arm spacing under OM (optical microscope) and the sites of precipitates under SEM.

Fig. 1.

Sampling method.

3. Results and Discussion

3.1. Solidification Microstructure of the Ingot

The solidification microstructure of the ingot is captured by a scanner and shown in Fig. 2. ‘V’ type segregation can be observed obviously in the upper area. 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.

Fig. 2.

Solidification microstructure of ingot.

The primary dendrite and secondary dendrite which grows from the primary one in position 1, 2 and 3 is shown in Figs. 3(a), 3(b) and 3(c) respectively. The measured value of secondary dendrite arm spacing in the corresponding site is 44.81 μm, 45.18 μm and 43.09 μm. The cooling rate during solidification is rough predicted by Eq. (1) which is proposed by Won and Thomas.¹⁹⁾ 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 calculated value in the position 1, 2 and 3 is 38.77°C/s, 37.90°C/s and 43.20°C/s respectively.   

$${\lambda }_{SDAS}(\mu \text{m})=143.9\cdot C_R{\hspace{0.17em}}^{-0.3616}\cdot C_C{\hspace{0.17em}}^{(0.5501-1.996C_C)}$$(1)

Fig. 3.

Secondary dendrite in (a) position 1, (b) position 2, (c) position 3.

3.2. Characteristics of Micron-sized Precipitates

A large number of (Ti,V) carbonitride can be observed in the samples. These precipitates could be further classified as the homogeneous Ti-rich ones and the inhomogeneous Ti-V-rich ones. As shown in Fig. 4, the homogeneous Ti-rich precipitate normally has a regular quadrangle shape and the size is mostly smaller than 5 μm. The black core in the precipitate is oxide the composition of which is mainly Al₂O₃.

Fig. 4.

Typical morphology of homogeneous Ti-rich (Ti,V) carbonitride.

The atomic contents of alloying elements tested by EDS in the different sites in Fig. 4 are listed in Table 2 and the corresponding ratios of Ti/V are added in the figure. A relatively low content of V is contained compared with that of Ti. Other elements, such as Mo and Cr are negligible. Element Fe may be more likely to originate from the iron matrix for that the content of Fe in the center is lower compared with that in edge. The face scanning mapping of elements Ti, V, C and N is shown in the right side of Fig. 4. There exists an obviously enriched region of N while that of C is ambiguous. It shows that Ti-rich precipitates are rich in N even though the content of light elements C and N could not be determined accurately by EDS analysis.

Table 2. Atomic contents of alloying elements in Ti-rich precipitates, at%.

Ti/V	Ti	V	Mo	Cr	Fe	C	N
5.85	39.24	6.71	0.68	0.68	4.02	15.65	33.02
6.14	34.30	5.59	0.61	0.74	4.50	20.39	33.86
6.04	36.20	5.99	0.47	0.57	1.26	17.64	37.86
5.96	34.33	5.76	0.64	0.71	5.93	16.53	36.80

Besides the homogeneous ones, a large number of inhomogeneous Ti-V-rich precipitates are also observed in the samples, as shown in Fig. 5. The numbers represent the atomic ratios of Ti/V in the sites the cross symbols refer to. From the dark gray area to the light gray area or the center to the edge in one precipitate, the ratio of Ti/V decreases gradually and even reaches zero at the end. Taking Fig. 5(d) for example, the face scanning mapping of elements Ti, V, Mo, Cr, C and N is shown in Fig. 6(a), and the line scanning mappings are shown in Figs. 6(b), 6(c), 6(d) and 6(e). The atomic contents of alloying elements in a part of points in Fig. 5(d) are listed in Table 3. It can be found that with the shift from the Ti-rich area to the V-rich one, the phase rich in nitrogen will transform the one rich in carbon. In other words, the precipitate precipitating first is rich in titanium and nitrogen, while the later one will be rich in vanadium and carbon. In addition, the contents of Mo and Cr increase gradually with the shift going.

Fig. 5.

Typical morphology of inhomogeneous Ti-V-rich (Ti,V) carbonitride, the number refers to the atomic ratio of Ti/V.

Fig. 6.

(a) Face scanning mapping and (b), (c), (d), (e) line scanning mapping of one typical Ti-V-rich (Ti,V) carbonitride.

Table 3. Atomic contents of alloying elements in Ti-V-rich precipitate, at%.

Ti/V	Ti	V	Mo	Cr	Fe	C	N
3.99	42.29	10.60	1.50	1.50	1.15	16.13	27.71
3.79	39.73	10.49	1.61	0.60	1.16	19.80	26.61
1.60	29.61	18.52	4.36	1.57	2.58	32.44	10.92
0.97	19.15	19.77	4.94	2.50	8.31	39.15	6.18
0.42	10.74	25.87	5.22	2.99	6.77	44.21	4.20
0.06	1.55	27.72	5.53	5.67	18.86	39.02	1.65

The atomic contents of Ti, V, Mo and Cr in all (Ti,V) carbonitrides observed in three samples are normalized and plotted into the ternary phase diagram of Ti−V−(Mo+Cr), as shown in Fig. 7(a). The ones in position 1, position 2 and position 3 are represented by half-solid square, half-solid triangle, and half-solid circle respectively. The compositional characteristics in three positions are consistent with each other even though the cooling rates are different. In addition, the atomic contents of Ti and V in (Ti,V) carbonitrides analyzed by EDS are plotted into a binary phase diagram and shown in Fig. 7(b). The ones in the homogeneous Ti-rich precipitates are represented by blue color. The ratios of Ti/V range from 3.36 to 9.53 and the arithmetic mean value is 6.35, while those in Ti-V-rich carbonitrides range from 0 to 7.60. The arrow refers to the composition evolving direction. As analyzed above, the precipitates precipitating later will be rich in vanadium.

Fig. 7.

(a) Normalized atomic ratios of Ti, V, Mo and Cr, and (b) atomic contents of Ti and V in (Ti,V) carbonitrides.

Mo-Cr-rich carbides can be also found, as shown in Fig. 8 with symbol ‘Mo-Cr’. The composition is listed in Table 4. A relatively low content of vanadium is contained. In addition, some sulfides also exist in the samples, singly or acting the nucleus of carbonitride, such as the black phase with symbol ‘S’ in Fig. 8. From the combining form of different types of phases, it can be speculated that Mo-Cr-rich carbides precipitate after (Ti,V) carbonitride and sulfide.

Fig. 8.

Combining form of different types of large precipitates.

Table 4. Atomic content of alloying elements in Mo-Cr-rich precipitate, at%.

V	Mo	Cr	Fe	C	N
8.96	17.66	16.67	9.55	45.38	1.78

The etching pattern of the samples under SEM is shown in Fig. 9. The gray area is the boundary of the dendritic crystals which grow toward the liquid steel during solidification. Alloying elements are enriched in the area due to the segregation during solidification. The black line is the grain boundary. Precipitates are marked according to the compositional characteristics. These precipitates are all located between the dendritic crystals. In other words, these precipitates are generated in the liquid steel during solidification.

Fig. 9.

Precipitating site of the large precipitates in (a) position 1 and (b) position 2.

3.3. Thermal Stability of the Primary Precipitates

The thermal stability of (Ti,V) carbonitride is assessed by quenching experiment. Samples with the size 10 mm×10 mm×10 mm were cut from the as-cast ingot and encapsulated in φ18 mm×88 mm quartz tubes. The tubes were filled with Ar gas so that the oxidation and decarburization were limited. The samples were austenitized at 1150°C and 1250°C for 0.5 h, 3 h and 6 h and immediately quenched in cold water. The results are shown in Figs. 10 and 11. The corresponding temperature and time are located on the top right corners. The precipitates are marked according to the dominant elements.

Fig. 10.

Morphologies of large phases after holding 0.5 h, 3 h and 6 h at 1150°C.

Fig. 11.

Morphologies of large phases after holding 0.5 h, 3 h and 6 h at 1250°C.

Ti-rich precipitates are stable at 1150°C and 1250°C. No obvious changes on the morphology occur, as shown in Figs. 10(a), 10(e) and 10(i), and 11(a), 11(e) and 11(i). The atomic contents of Ti and V in Ti-rich precipitates after different holding times are shown in Fig. 12. The original values are represented by the black half-solid square, and the ones after 0.5 h, 3 h and 6 h are represented by blue half-solid circle, red half-solid circle and green half-solid circle respectively. The changes on the atomic ratios of Ti/V are also little even though some shift of the total values toward the higher value region exist.

Fig. 12.

Atomic contents of Ti and V in (Ti,V) carbonitrides after different holding times at (a) 1150°C and (b) 1250°C.

For inhomogeneous Ti-V-rich precipitates, the ones with low ratios of Ti/V start to decompose even after 0.5 h at 1150°C, especially the one with elongated shape, as shown in Fig. 10(c). The numbers represent the atomic ratios of Ti/V. The precipitates with high contents of titanium are relatively stable. However, the decomposition is also observed after holding 3 h at 1150°C, as shown in Fig. 10(f). After 6 h at 1150°C, the decomposition goes on and a part of ones with low ratio of Ti/V has transformed into a distribution of small precipitates, as shown in Figs. 10(k) and 10(l). The transformation is more obvious at 1250°C. The decomposition can be observed even after 0.5 h, as shown in Figs. 11(b), 11(c) and 11(d). The precipitates left are the ones or the parts with higher ratio of Ti/V, as the thermal stability of Ti-rich ones analyzed above. The thermal stability of (Ti,V) carbonitride is similar to the study of M. Prikryl.¹¹⁾ He thought that the thermodynamic driving force combining with the minimization of the interfacial energy would contribute the dissolution and shape transformation. Then the thermal stability of the precipitates analyzed above will be discussed below based on the equilibrium precipitation.

Mo-Cr-rich carbides are unstable under high temperature. They start to dissolve even after 0.5 h at 1150°C. These precipitates could be not observed after 3 h at 1150°C and even 0.5 h at 1250°C.

4. Analysis and Discussion

4.1. Equilibrium Precipitation

The equilibrium precipitation was calculated by Thermo-Calc thermodynamic software. Database TCFE7 was applied and miscibility gap in (Ti,V)(C,N,Vacancy) was considered. The total system was set as 1 g. The calculated result between 1800–1000°C is shown in Fig. 13. Eight types of phases exist in the equilibrium diagram, that is, liquid steel, ferrite, austenite, oxide (Al,Ti)O, Ti₄C₂S₂, sulfide MnS, FCC_A1#2 and FCC_A1#3. FCC_A1 is a FCC phase with disordered structure type A1 defined in the database in Thermo-Calc.

Fig. 13.

Equilibrium precipitation.

The liquidus temperature is about 1470°C. The final composition of liquid steel before solidification is listed in Table 5. The soluble aluminum [Al] left in liquid steel is 0.00561 wt% for the generation of oxide (Al,Ti)O and is consistent with the tested composition listed in Table 1. It shows that (Al,Ti)O is mainly generated in the liquid steel before solidification. The atomic content of Al, Ti and O in (Al,Ti)O is shown in Fig. 14(a). The content of Ti is low. The composition of the oxide is also consistent with that of the oxide core in Fig. 4. The content of Ti left in liquid steel [Ti] is 0.0338 wt%. Except for Ti and Al, the contents of other elements are unchanged.

Table 5. Final composition of liquid steel before solidification, wt%.

C	Si	Mn	Cr	Mo	V	[Ti]	[Al]	S	O	N
0.35	1.03	0.40	5.29	1.51	0.97	0.0338	0.00561	0.0059	0.000221	0.0033

Fig. 14.

Atomic content of alloying elements in (a) (Al,Ti)O, (b) FCC_A1#2 and (c) FCC_A1#3.

The atomic content of alloying elements in FCC_A1#2 and FCC_A1#3 is shown in Figs. 14(b) and 14(c) respectively. FCC_A1#2 is a Ti(C,N) phase. Other elements, such as V, Mo and Cr are negligible. With the decrease of temperature, the ratio of N decreases while that of C increases. However, the tendency is opposite when the temperature drops below 1134°C for the precipitation of FCC_A1#3. FCC_A1#3 is a V-rich carbide. A relatively low content of Ti is contained and reaches zero with the decrease of temperature. The content of Mo and Cr is negligible.

The equilibrium carbonitride at 1150°C and 1250°C is Ti(C,N). The primary precipitates observed in the samples are not equilibrium ones. According to the analysis of M. Prikryl, the thermodynamic driving force will contribute to the dissolution with the homogeneous of composition in iron matrix. The driving force will be larger for the precipitates with low content of titanium. Combining with the concentration gradient of vanadium in the Ti-V-rich precipitates and the larger diffusion coefficient of vanadium in the iron matrix,¹¹⁾ the dissolution accompanying with the inner decomposition will take place first in the region with low ratio of Ti/V under high temperature. In addition, the minimization of the interfacial energy of the precipitates will further contribute to the transformation, especially the ones with elongated shape, as shown in Figs. 10(g) and 10(h). Based on the equilibrium precipitation, the homogeneous Ti-rich precipitates observed in the ingot might be also unstable. However, it will need more time to finish the transformation.

4.2. Thermodynamic Calculation for Precipitating Process during Solidification

During solidification, microsegregation will occur with the development of solidification structure. Taking position 1 and 3 for example, with the development of columnar crystal, the solutes are gradually rich in liquid steel between the secondary dendrite arms. The primary precipitates will precipitate in the liquid steel when the contents of alloying elements reach the precipitating condition.

The solidification process in position 1 and 3 can be speculated through the Scheil-Gulliver model in Thermo-Calc 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 alloying elements are homogeneous in liquid steel, 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. The difference between the calculated solidification process and the actual one could be further reduced by the remove of primary carbonitrides in each calculation step. The composition in Table 5 is used as the initial composition of liquid steel during the step calculation. The content of O is ignored, or some unreasonable oxide will appear in Scheil Model.

The calculation result is shown in Fig. 15. There is no value at solid fraction 1.0 for that the final solute concentration in liquid steel will become infinite and the value is insignificant in Scheil model. Six phases will appear during solidification, that is, the matrix ferrite, the matrix austenite, and the primary precipitates FCC_A1#2, MnS, M₇C₃ and M₆C except for the oxide (Al,Ti)O which is mainly generated in the liquid steel before solidification. M₇C₃ and M₆C are carbides and defined in the database in Thermo-Calc according to the stoichiometric formula. During solidification, the primary precipitate FCC_A1#2 first appears and then MnS, M₇C₃ and M₆C precipitate successively. The corresponding solid fractions of precipitation for the primary phases are listed in Table 6. The precipitation sequence is consistent with the actual one shown in Fig. 8.

Fig. 15.

Solidification process and the primary precipitates.

Table 6. Solid fraction of precipitation for the primary phases.

Primary precipitates	FCC_A1#2	MnS	M7C3	M6C
Solid fraction	0.727	0.946	0.964	0.972

The atomic content of alloying elements in FCC_A1#2, M₇C₃ and M₆C is shown in Figs. 16(a), 16(c) and 16(d) respectively. Figure 16(b) is the sectional representation for the dotted box region in Fig. 16(a). FCC_A1#2 phase is a (Ti,V)(C,N) type precipitates. The one precipitating first is rich in Ti and N. A certain content of C is contained and increases with the development of solidification. When the solid fraction reaches 0.908, the contents of vanadium and carbon in the newly generated precipitates increase sharply, and V-C-rich carbonitride with a certain content of Ti will precipitate. A certain content of Mo and Cr are contained in the V-C-rich phase and increase gradually.

Fig. 16.

Atomic content of alloying elements in (a) FCC_A1#2, (b) the sectional representation of FCC_A1#2, (c) M₇C₃, (d) M₆C.

The compositional characteristic of FCC_A1#2 seems to be consistent with that of (Ti,V) carbonitrides observed in the samples. To compare with the experimental results in Fig. 7(a), the atomic contents of Ti, V, Mo and Cr in FCC_A1#2 are normalized and plotted into the Ti−V−(Mo+Cr) ternary phase diagram, as shown in Fig. 17. The empty triangle represents the calculated results and the curve with an arrow represents the compositional evolving direction with the development of solidification. The compositional distribution of (Ti,V) carbonitrides in Fig. 7(a) is also added into the diagram and presented as the half-solid circle. The calculated results agree with the experimental results well. In addition, the calculated results show that Ti(C,N) precipitates first in the liquid steel during solidification. However, no pure Ti(C,N) is observed in the samples. This may be the diffusion result of vanadium from iron matrix during the subsequent cooling process.

Fig. 17.

Compositional comparison between the calculated results and experimental results for (Ti,V) carbonitride.

M₇C₃ is a Cr-Fe-rich carbide and has a little content of V and Mo, while M₆C is a Mo-Fe-rich carbide with a certain content of Si and little content of V and Cr. As analyzed in Ref. 20) which studied the characteristics of primary precipitates in Nb-Ti-microalloyed H13 ingot, Mo-Cr-rich carbide observed in the samples may be the mixture of M₇C₃ and M₆C.

Based on the analysis above, the generating mechanism of the primary precipitates in position 1 and 3 could be described in Fig. 18. During solidification, the solutes C, N, Ti, V, Mo and Cr are rejected into the liquid steel with the development of solidification structure. Ti-N-rich carbonitride precipitates first when the nucleation condition has achieved, and the composition changes little until the generation of Ti-V-rich precipitates. Ti-V-rich carbonitride will precipitate singly or on the Ti-rich phase which has already existed in the liquid steel. With the development of solidification, the contents of Ti and N in the newly generated Ti-V-rich precipitates decrease and reaches zero eventually, while that of V and C increases, resulting in the formation of inhomogeneous Ti-V-rich precipitates. Sulfide will appear at the end of solidification, acting the nucleus of Ti-V-rich carbonitride or Mo-Cr-rich carbide which precipitates latest.

Fig. 18.

Generating process of primary precipitates during solidification.

As analyzed in Section 3.2, the compositional characteristics of (Ti,V) carbonitrides in different positions in the ingot are consistent with each other. It may be due to the low diffusion amount of the carbonitride forming elements in solid steel during solidification under the high cooling rate. Actually, the generating mechanism of primary precipitates in position 2 with ‘V’ type segregation is different from position 1 and 3.²⁰⁾ However, it seems 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 detriment of the Ti-V-rich carbonitride could be reduced to some extent by heat treatment for the dissolution or decomposition of precipitates with low ratio of Ti/V. However, it is useless for the one with high content of Ti. Furthermore, heat treatment under relatively lower temperature will need more time to dissolve the large precipitates, while that under higher temperature may result in the grain coarsening of austenite and even local remelting of iron matrix. Adjusting the content of Ti and N, or taking reasonable measures to refine as-cast structure^21,22,23,24) is more direct and effective.

5. Conclusions

(1) A large number of (Ti,V) carbonitrides exist in H13+0.034 wt% Ti ingot, including the homogeneous Ti-rich ones and the inhomogeneous Ti-V-rich ones. The atomic ratios of Ti/V in Ti-rich precipitates range from 3.36 to 9.53, while those in Ti-V-rich precipitates range from 0 to 7.60. The sizes of Ti-rich ones are normally smaller than 5 μm. In addition, a certain number of sulfide and Mo-Cr-rich carbide exist in the ingot.

(2) Ti-rich carbonitride is stable at 1150°Cand 1250°C and no obvious change on shape and composition is observed. Ti-V-rich precipitates are unstable. The ones with low ratio of Ti/V start to decompose after holding a short time, while the transformation of the ones with higher ratio of Ti/V will need more time. No Mo-Cr-rich carbide is left after 3 h at 1150°C or even 0.5 h at 1250°C.

(3) The large precipitates are generated during solidification. The generating process could be well speculated by Thermo-Calc. During solidification, Ti-N-rich carbonitride precipitates first and the composition changes little until the generation of Ti-V-rich precipitates. With the development of solidification, the contents of Ti and N in newly generated Ti-V-rich precipitates decrease, while that of V and C increase, resulting in the formation of inhomogeneous Ti-V-rich precipitates. Sulfide and Mo-Cr-rich carbides precipitate at the end of solidification.

Acknowledgments

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

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