2017 Volume 57 Issue 7 Pages 1252-1260
In this study, different amounts of TaC powders (10, 15 and 20 mass%) are mixed and added to Vanadis 4 tool steel powder. The composite powders are sintered at 1250, 1300 and 1325°C for 1 h, respectively. After that, a series of heat treatments is performed to determine the optimal parameters of the Vanadis 4 composites. The experimental results show the optimal vacuum sintering temperature for the Vanadis 4 composites to be 1300°C. The Vanadis 4 specimens with a 10% TaC added possess the highest transverse rupture strength (TRS) value of 1974.2 MPa and a suitable hardness (82.6 HRA). When the optimally sintered specimens of the Vanadis 4 composites then undergo a heat treatment, the TRS shows a significant increase to 2069.4 MPa, while the hardness declines slightly to 81.5 HRA. In addition, the microstructural evaluation reveals that the plate-shaped carbides (M7C3) located on the grain boundaries disappear after the TaC powders are added and that clustered TaC carbides appear in the grain boundaries instead. Meanwhile, with TaC as the nucleation sites of the VC carbides, the result is the generating of VC carbides in the grain boundaries. Moreover, the VC carbides decompose and re-precipitate the refined carbides (Fe3C), which are uniformly distributed within the grains. The results also show that heat treatment effectively improves the microstructure and strengthens the Vanadis 4 composites.
Vanadis 4 is a high vanadium and chromium tool steel, which is widely used for cold extrusion tooling, fine blanking and IC molding, due to its high wear resistance, toughness and suitable hardness. Furthermore, the primary strengthening mechanisms include solid-solution strengthening and precipitation-hardening by chromium and vanadium carbides, respectively. In order to further improve the tool life, over the past 10 years, industrial tool steels for cutting or wear of materials have used metal-ceramic composite materials, whereas, powder metallurgy (P/M) methods offer two very different types of materials with a view to achieving lightweight, high strength, hardness and wear resistance tool steels.1)
Up to now, metal matrix composites (MMCs) have been widely studied and applied in infrastructure industries and aerospace technology because of their capability to provide functional properties that include excellent wear resistance, high ductility and attractive thermal characteristics.2,3) Additionally, many investigations have revealed that reinforced MC carbide (TaC, VC, NbC and TiC) in ferrous metal matrix will efficiently enhance the mechanical performance.4,5,6) Among them, tantalum carbide (TaC) has excellent hardness (1800 HV50 kgf·mm−2), high melting temperatures (3422°C) and good chemical and thermal stability with Fe-based alloys; therefore, it shall be a great option for reinforcing in Vanadis 4 tool steels.7)
One of the main advantages of P/M procedure is its flexibility to prepare different powder mixtures admixing different hard and/or soft particulate phases to the basic high-alloyed steel powder. In this way, special composite or gradient materials with target properties for a precisely defined application (e.g. improved wear resistance, higher thermal or corrosion resistance, better machinability and selflubricity, etc.) are synthesized.8,9,10)
Generally speaking, in higher carbon content alloyed steels, the martensite finish temperature (Mf) is below 0°C, which means that at the end of the heat treatment, a low percentage of austenite may be retained at room temperature. The retained austenite as a soft phase in steels could reduce the product life. In order to resolve the problems mentioned above, the cryogenic (also known as sub-zero) treatment is used to transform the retained austenite into martensite. Thus, the retained austenite is reduced and higher wear resistance is obtained. The cryogenic treatment can be as an effective way to improve the properties of tool steels, which is widely used for high precision parts and components since it enhances the transformation of austenite to martensite.11,12)
Previous studies have indicated that Vanadis 4 tool steel is a P/M and cold work tool steel which offers an extremely good combination of wear resistance and ductility. After commercial heat treatments, a hardness of 81.2 HRA (60 HRC) can be achieved.13) Moreover, TaC is one of the best grain growth inhibitors. Morton et al. showed that TaC, unlike VC, is unaffected by the carbon content of the WC used to control the grain growth process in WC-Co composites.14,15) The question of whether or not different amounts of TaC particles added to Vanadis 4 are effective as a strengthening phase, in order to improve the mechanical properties, as well as the distribution and size of the carbides for Vanadis 4 tool steels, remains unanswered. The aim of this work was to explore a series of vacuum sintering and heat treatments for Vanadis 4 steels in order to examine the effects on the microstructure and mechanical properties of TaC-strengthened Vanadis 4 steels.
The aim of this study was to use Vanadis 4 tool steel powders and add different ratios of TaC powders as a strengthening phase in order to explore the effects of a series of vacuum sintering and heat treatments. In this work, gas-atomized Vanadis 4 tool steel powders were used as a substrate. The morphology of Vanadis 4 powders shows their having near round and equiaxed particle characteristics. In order to obtain more uniform powders, particle sizes smaller than 38 μm were sifted from the original powders. As a result, the mean particle size of the Vanadis 4 tool steel powders was 21.5±0.1 μm, as shown in Fig. 1(a).
The SEM images of the surface morphology of (a) the sifted Vanadis 4 steel powders, (b) the original TaC powders, and (c) Vanadis 4-TaC powders after ball mixing for 6 h.
The chemical compositions (mass%) of the Vanadis 4 tool steel powders are listed in the Table 1. Furthermore, the surface morphology of TaC powders is shown in Fig. 1(b). The mean particle size of the TaC powders was about 1.3±0.1 μm; moreover, the shape of the powders was a polygon and irregular, and there were no smooth or undulating surfaces. The different amounts of TaC powders (10, 15 and 20 mass%) were mixed and added to Vanadis 4 tool steel powders, designated as T10, T15 and T20, hereafter. The mixed powders were milled by using the WC balls for 6 h. Figure 1(c) shows the morphology under the effect of mechanical alloying by ball milling at 300 rpm for 6 h, and the mixed alloy powders produced a significantly plastic deformation. After milling, the PVA (polyvinyl alcohol) as a binder was added. The green body (40 × 6 × 6 mm3) of the powder specimen was produced under a uniaxial pressure at 350 MPa for 300 s.
| C | Si | Mn | Cr | Mo | V | Fe | |
|---|---|---|---|---|---|---|---|
| Vanadis 4 | 1.51 | 1.0 | 0.4 | 8.0 | 1.5 | 4.0 | Bal. |
In this study, the vacuum sintering was conducted at 1250, 1300 and 1325°C for 1 h in a 5 × 10−3 Pa, respectively. When the sintering temperature was increased to 1350°C, the melting phenomenon appeared in the sintered specimens; thus, the sintering temperature was lowered to 1325°C. In addition, a series of heat treatments (quenching followed by sub-zero and tempering) was performed, by which the samples were heated to 1020°C and maintained at that temperature for 100 min for quenching, with 0.5 MPa of N2 used as the quenching media; the samples were then subjected to sub-zero treatment at a temperature of −150°C for 60 min. Simultaneously, the tempering temperature was held at 480°C for 180 min and repeated two times.
To evaluate the sintered behavior of Vanadis 4 tool steels added TaC powders by vacuum sintering and heat treatments, the volume shrinkage, apparent porosity, hardness, transverse rupture strength (TRS) tests and microstructure observations were performed. Microstructural observations of the specimens were performed by optical microscopy (OM) and scanning electron microscopy (SEM, Hitachi-S4700). Volume shrinkage and porosity test were conducted in accordance to the ASTM C830 standard. Hardness tests were performed by Rockwell A hardness (HRA, Indentec 8150LK) with a loading of 588.4 N, which followed the ASTM B294 standard. The Hung Ta universal material test machine (HT-9501A) with a maximum load of 245 kN was used for the TRS tests (ASTMB528-05). Meanwhile, the TRS was obtained by the equation Rbm = 3FLk/2bh2, where Rbm is the TRS, which is determined as the fracture stress in the surface zone, F is maximum fracture load. In this work, L was 30 mm, k was chamfer correction factor (normally 1.00–1.02), b and h were 5 mm, respectively. The specimen dimensions of the TRS test were 5 × 5 × 40 mm3 and tests at least three pieces.
Figure 2 shows the volume shrinkage rate and porosity as a result of the different sintering temperatures for the various amounts of TaC added to the Vanadis 4 specimens. As Fig. 2(a) shows, the volume shrinkage rate of the T10, T15 and T20 specimens had a similar trend, i.e., that when the sintering temperature reached 1250°C, all specimens had a relatively low volume shrinkage rate resulting from the restricted diffusion of the TaC particles. As the sintering temperature increased to 1300°C, the volume shrinkage rate (37.4%) of the T10 specimen showed a significant increase, as compared with the optimally sintered Vanadis 4 specimens (29.9%). Furthermore, when the sintering temperature was further increased to 1325°C, the volume shrinkage rate of all specimens was over 37%. It is reasonable to suggest that it was easier for a small amount of added TaC powders to be rearranged and so enter the interspaces of the Vanadis 4 particles during the sintering process. As a result, the high density of the sintering materials was acquired. As for the porosity level, as shown in Fig. 2(b), the porosity of the T10 and T15 specimens was less than 1% after sintering at 1300°C for 1 h (> 98% theoretical density), and especially that of the T10 specimen (porosity of 0.1%). Actually, an excess of TaC powders obstructed the diffusion of the liquid phase during the sintering process. In this case, the porosity of the T20 specimen decreased to 6.3% after vacuum sintering at 1300°C for 1 h as a result of an incomplete liquid phase sintering (LPS). Notably, the porosity level of all specimens decreased to less than 1% after vacuum sintering at 1325°C for 1 h. The near-full theoretical sintered densities of all sintered specimens were acquired.
Comparison of the volume shrinkage and apparent porosity of various mass% TaC added Vanadis 4 by the different sintering temperatures.
As the comparison of Figs. 2(a) and 2(b) shows, the specimens with higher amounts of TaC powders added required a higher temperature to provide sufficient sintering energy for a complete LPS. Clearly, the T10 specimen had a higher volume shrinkage rate than the T15 or T20 specimens after sintering at 1300°C. When the sintering temperature was raised to 1325°C, the volume shrinkage rate of all specimens rapidly increased to over 37%. However, the volume shrinkage rate of the T10 specimen showed only a slight variation as the sintering temperature increased to 1325°C. The results suggested that a sintering temperature of 1300°C corresponded to the LPS temperature for the T10 sintered specimen. As for the porosity, that of the T20 specimen was relatively high after sintering at 1300°C; however, it was decreased to 0.24% by the draining off the pores from the high-energy grain boundaries when the sintering temperature was raised to 1325°C. As a result, the near-full density of the Vanadis 4 composite materials was achieved.
Figures 3(a)–3(c) show the OM images of the T10, T15 and T20 specimens after sintering at 1250°C for 1 h. The amount and size of the residual porosity increased as the added amount of TaC increased. The TaC powders distributed and formed clusters of TaC carbides on the grain boundaries, and these increased as the added amount of TaC increased. As compared with the OM image of the original Vanadis 4 specimens (Fig. 3(d)), two kinds of carbides appeared in the specimens: numerous spherical-shaped carbides (MC) within the grains, and some plate-shaped carbides (M7C3) appeared on the grain boundaries. As a result, the sintering temperature at 1250°C for the Vanadis 4 composite materials was not high enough. When the sintering temperature was increased to 1300°C, the OM images of the T10, T15 and T20 specimens displayed more distinct microstructures, as shown in Fig. 4. As Figs. 4(a) and 4(b) show, some small pores still existed within the matrixes after sintering at 1300°C for 1 h. Nevertheless, the pores showed an obvious decrease, as compared with the composite specimens sintered at 1250°C (Fig. 3). On the other hand, increasing the amount of TaC powders added caused a significant decrease in the mean grain size (Vanadis 4 at 33.1 μm, T10 at 16.5 μm, T15 at 13.0 μm, respectively). This result indicated that increasing the amount of TaC powders added helped restrain grain growth. Notably, the T20 specimen did not attain an ideal sintered microstructure after sintering at 1300°C for 1 h, as shown in Fig. 4(c). In addition, the carbide distribution was significantly different, as compared with the Vanadis 4 specimen (Fig. 3(d)). The plate-shaped carbides in the grain boundaries obviously disappeared, and the clustered TaC carbides appeared on the grain boundaries in their place. It is reasonable to suggest that the TaC particles distributed on the grain boundaries inhibited the grain growth during the sintering process. However, the TaC particles underwent high-temperature diffusion sintering and formed continuous clusters of TaC carbides on the grain boundaries. An excessive clustering of TaC carbides around the grain boundaries was disadvantageous to the TRS.
The OM images of various mass% TaC added Vanadis 4 after sintering at 1250°C/1 h. (a) T10, (b) T15, (c) T20, and (d) original Vanadis 4 specimen (non TaC added).
The OM images of various mass% TaC added Vanadis 4 after sintering at 1300°C/1 h. (a) T10, (b) T15, and (c) T20.
Figure 5 shows the OM images of the T10, T15 and T20 specimens after sintering at 1325°C for 1 h. A further comparison with Fig. 4 shows that the average grain size obviously increased as the sintering temperature increased (T10 from 16.5 to 20.8 μm, T15 from 13.0 to 19.2 μm, respectively). Although the T20 specimen still had some pores within the matrix, as shown in Fig. 5(c), it possessed a more distinct microstructure after sintering at 1325°C for 1 h. However, all the composite specimens nearly melted after sintering at 1350°C for 1 h, which rendered the specimens unsuitable for subsequent research and testing. Perhaps when fabricating Vanadis 4 composites with a large number of strengthening phases, other sintering technologies, such as hot pressing or HIP (hot isostatic pressing), must be utilized to improve the microstructure.
The OM images of various mass% TaC added Vanadis 4 after sintering at 1325°C/1 h. (a) T10, (b) T15, and (c) T20.
In order to further investigate the morphology and compositions of precipitates of Figs. 3, 4, 5. This study utilized the SEM and EDS’s analysis, as shown in Fig. 6. As to the carbide distribution, our previous study of Vanadis 4 tool steel confirmed that both the plate-shaped carbides in the grain boundaries and the spherical-shaped carbides within the grains were V-rich and Cr-rich M7C3 and MC carbides, respectively.13) Moreover, all the Vanadis 4 composite materials had a similar morphology and compositions of carbide. Taking the specimen (T10) sintered at 1300°C for 1 h as an example, as seen in Fig. 6(a), a few small spherical-shaped carbides within the grains can be observed. The EDS analysis results revealed that the spherical-shaped carbides within the grains (location 1) were V-rich VC carbides (28.62 at% V); the compositions are listed in Fig. 6(a)-1. Besides, the clustered TaC carbides in the grain boundaries appeared to replace the plate-shaped carbides. The EDS analysis results showed that a high content of V-rich (14.66 at%) appeared in the TaC carbides in spite of Ta-rich, Fe-rich and C-rich, as listed in Fig. 6(a)-2. It is possible to say that the TaC served as the nucleation sites of the VC carbides, which resulted in the generating of VC carbides in the grain boundaries and the decrease in VC carbides within the grains (Fig. 4(b)).
The SEM images and EDS analysis location of the optimal sintered T10 and T10+HT specimens (a) SEM images of T10, (a)-1 and (a)-2 are EDS results of T10, and (b) SEM images of T10+HT, (b)-1 and (b)-2 are EDS results of T10+HT.
Figure 7 shows the hardness and TRS test results of the T10, T15 and T20 specimens at different sintering temperatures. All the completed values of hardness and TRS are shown in Tables 2 and 3. As Fig. 7(a) shows, the hardness of all specimens was relative to the porosity level (Fig. 2(b)) after sintering at the different temperatures. The hardness was relatively low when the sintering temperature was 1250°C for 1 h, indicating that the specimens did not reach a complete LPS. However, the T10 and T15 specimens increased to an ideal value (> 82 HRA) after the sintering temperature reached a complete LPS condition (1300°C for 1 h). Furthermore, the hardness did not improve when the sintering temperature was raised to 1325°C. These results were further compared with Table 2. All specimens having reached the complete LPS possessed a relatively high hardness (> 82 HRA). The hardness was higher than that of the Vanadis 4 tool steel (80.5 HRA), which improved only slightly with increases in the amount of TaC added. It was therefore surmised that reinforced TaC powders added to the Vanadis 4 steel matrix effectively enhanced the hardness. Moreover, the Vanadis 4 composites reached a complete LPS condition, and the amount of TaC added to Vanadis 4 and the evolution of sintering temperatures did not seem to effect a variation in hardness.
Comparison of the hardness and TRS of various mass% TaC added Vanadis 4 by the different sintering temperatures (a) hardness, and (b) TRS.
| Sintering Temperatures (°C) | Vanadis 4 | T10 | T15 | T20 | T10+HT |
|---|---|---|---|---|---|
| 1250 | 80.5 | 43.7 | 49.2 | 43.7 | – |
| 1300 | – | 82.6 | 82.7 | 78.2 | 81.6 |
| 1325 | – | 82.1 | 82.3 | 82.3 | – |
| Sintering Temperatures (°C) | Vanadis 4 | T10 | T15 | T20 | T10+HT |
|---|---|---|---|---|---|
| 1250 | 1262.2 | 88.6 | 229.8 | 145.1 | – |
| 1300 | – | 1974.2 | 1817.3 | 944.4 | 2069.4 |
| 1325 | – | 1723.7 | 1720.3 | 1634.8 | – |
As Fig. 7(b) shows, the TRS value of the T10 and T15 specimens first increased and then decreased. The T20 specimen tended to increase as the sintering temperature increased from 1250 to 1325°C. The T10 and T15 specimens possessed a better TRS after sintering at 1300°C, with the T10 specimen having the highest TRS value (1974.2 MPa). Our previous study indicated the precipitated carbides and porosity levels of the Vanadis 4 composite materials to be important factors directly affecting the TRS.16) If the sintered specimen were not fully densified, the materials would possess numerous small pores. When the specimen was subjected to extra stress during the TRS tests, cracks grew along the internal pores, resulting in rapid fracturing. The results when further compared with the Vanadis 4 specimen, as also shown in Fig. 7(b), show that the TRS value for T10 specimen obviously improved after sintering at 1300°C for 1 h. It is possible to then say that a suitable amount of carbides in the grain boundaries and the lower porosity (0.1%) effectively improved the TRS. After the sintering temperature was increased to 1325°C, the TRS value of the T20 specimen dramatically increased, while that of the T10 and T15 specimens slightly decreased. However, the T20 specimen still had a slightly high porosity level, which was disadvantageous to the TRS. On the other hand, although the porosity of the T10 and T15 specimens decreased by less than 0.2%, the high sintering temperature (1325°C) caused a grain-coarsening phenomenon which seemed to significantly affect the TRS values. Simultaneously, the higher amounts of TaC in the T15 specimen was more effective in inhibiting the grain-coarsening, so the TRS decreased more slowly.
Figure 8 shows the fracture feature of the T10, T15 and T20 specimens after sintering at 1300°C and 1325°C, respectively. The fracture feature of the T10 specimen was similar to that of the T15 specimen after sintering at 1300°C for 1 h. Many small cleavages could be observed, as shown in Figs. 8(a) and 8(b). The presence of a greater number of dislocations was generated by the plastic extension when the fracture was under continually increasing loads. In fact, the carbides resisted the movement of dislocation, and then the strength improved.17) As the load increased, the general aspect of the fracture showed the dimpled ruptures and small carbides cleavages resulting from the ductile behavior. Thus, there were many obvious dimple ruptures appearing on the fracture surface (as shown by an arrow). In this study, the TaC was the hard and brittle reinforced carbides in the grain boundaries contributed to the brittleness. The fracture surface observation was in agreement with the finding that the small cleavages were easily generated on the grain boundaries. As a result, the main fracture mode was an intergranular fracture. Particularly, the T20 specimen still had many pores between the TaC particles in the grain boundaries, which soon lead to cracks, indicating that the density of the sintered specimens was insufficient after sintering at 1300°C, as shown in Fig. 8(c). Increasing the sintering temperature (1325°C) can help increase the density of sintered specimens, but it leads to the grain-coarsening phenomenon. As Fig. 8(d) shows, the fracture feature of the T20 specimen was similar to that of the T10 and T15 specimens (sintered at 1300°C). Generally, grain coarsening is disadvantageous to the TRS. From the above results and discussion, it was possible to speculate that the T10 specimen had the optimal TRS value and a suitable hardness after sintering at 1300°C for 1 h. The heat treatment is discussed below.
The fracture surfaces of various mass% TaC added Vanadis 4 (a) 1300°C sintered-T10, (b) 1300°C sintered-T15, (c) 1300°C sintered-T20, and (d) 1325°C sintered-T20 specimens.
Figure 6(b) shows that SEM images of the optimal T10 specimen after the heat treatment. As compared with Fig. 6(a), the T10 plus heat-treated (T10+HT) specimen clearly had a reduced number of carbides in the grain boundaries. Meanwhile, a large number of refined carbides were generated within the grains. A reasonable explanation of this effect was that the VC carbides dissolved into the matrix, while a large number of refined carbides within the grains re-precipitated. In this study, the clustered carbides (TaC) formed in the grain boundaries during the sintering process, and the heat treatment could not change the distribution of the TaC particles, only control the shape and distribution of the VC carbides. Thus, TaC powders added to Vanadis 4 as a strengthening phase made them pile up and cluster in the grain boundaries (Fig. 6(b)). A magnification of the SEM image shows the spherical-shaped carbides to be randomly dispersed, while the refined carbides are uniformly distributed within the grains, as also shown in Fig. 6(b). The EDS analysis of the spherical-shaped carbides and refined carbides, which located in the 1 and 2 regions, revealed VC carbides and Fe3C carbides, respectively. Their compositions are listed in Figs. 6(b)-1 and 6(b)-2. The results confirmed that the VC carbides decomposed and re-precipitated as spherical-shaped VC and refined Fe3C carbides within the grains after the heat treatment. Both carbides (VC and Fe3C) precipitating within the grains would be advantageous to the TRS as a result of the dispersion strength. In addition, some small pores still existed among the high-temperature diffusion sintered TaC carbides after the heat treatment.
As mentioned previously, it is well known that sub-zero treatment effectively increases the hardness by decreasing the retained austenite. In the present study, the hardness of the T10 sample was about 82.6 HRA, whereas the hardness of the quenched, sub-zero treated and double tempering sample was about 81.6 HRA (Table 2). It is possible to say that the high hardness of the T10 sample was owing to the martensite morphology and the dissolved alloying elements in the martensite and austenite. In other words, the T10 specimen was not tempered by the martensite. The increase in the martensite volume fraction was most likely responsible; however, sub-zero treatment could have encouraged carbide formation during tempering, which led to the increased hardness.18) Indeed, a previous study indicated that in the sub-zero treatment, the variation in hardness was dominated by the tempering process, which released the internal stress from the tempering treatment and resulted in the decrease in hardness.19) The hardness test result was in agreement with the previous study in that the slight decrease in hardness could be ascribed to the release of the internal stress.
As for the TRS, the T10+HT specimen revealed a significant increase, as listed in Table 3. The TRS value obviously increased to 2069.4 MPa. It was reasonable to infer that the dissolving of the VC carbides in the grain-boundaries and the VC and Fe3C carbides re-precipitating within the grains resulted in the increase in the TRS. In the TRS test, the refined carbides resisted the movement of dislocation, and then the strength improved as the load increased. The fracture surface of such a ductile fracture appeared dimpled when observed on SEM images, as shown in Fig. 9. Ductile fractures occur through the formation and coalescence of micro-voids (dimples) along the fracture path. The fracture surface observation also found many small voids, which would have been from the breaking away of carbides which left the voids. Although heat treatment did not improve the hardness of the Vanadis 4 composite materials, it was indicated that the T10 specimen sintered at 1300°C for 1 h, followed by a heat treatment, possessed the optimal mechanical properties.
The fracture surfaces of the 1300°C sintered-T10 specimen after heat treatment.
In this study, the T10 specimen possessed the highest TRS value (1974.2 MPa) and the hardness reached 82.6 HRA after sintering at 1300°C for 1 h. The TRS and hardness of the T10 specimen showed significant improvement as compared to the Vanadis 4 tool steel. Also, the mechanical property of the composite materials dramatically increased after adding 10% TaC powder to the Vanadis 4 tool steel.
For the T10 specimen, the spherical-shaped VC carbides uniformly dispersed within the grains and the clustered TaC and VC carbides appeared on the grain boundaries after the optimal sintering temperatures were applied. When the specimens underwent a series of heat treatments, the VC carbides decomposed and re-precipitated as spherical-shaped VC and refined Fe3C carbides within the grains, which helped improve the strength.
Sub-zero and heat treatments effectively improved the distribution and size of the VC carbides and the TRS value of the Vanadis 4 composite materials, while the hardness displayed a slight decrease. The decrease in hardness could be ascribed to the release of internal stress after the tempering process. In this work, sub-zero and heat treatments did not change the distribution of the TaC particles, but controlled the shape and distribution of the VC and Fe3C carbides. As a result, the optimal TRS (2069.4 MPa) and suitable hardness (81.6 HRA) were obtained for Vanadis 4 tool steel by adding 10% TaC after sintering at 1300°C for 1 h, as well as sub-zero and heat treatments.
This research is supported by the Ministry of Science and Technology of the Republic of China under Grant No. MOST 105-2221-E-027-021-. The authors would like to express their appreciations for ASSAB STEELS TAIWAN CO., LTD.
