2024 Volume 64 Issue 1 Pages 154-164
This research added different ratios of Ta50Nb50C and B4C powders to Vanadis 60 high-speed steel powders, and then, sintered the Vanadis 60 composite from 1205 to 1225°C by vacuum sintering for 1 hour. Subsequently, a series of heat treatments were conducted, including quenching, sub-zero, and tempering processes. The study results show that the optimal sintering temperature and the additive ratio for Vanadis 60 composites were 1205°C and 0.975 mass% Ta50Nb50C-0.025 mass% B4C powders, respectively. Simultaneously, the transverse rupture strength (TRS) value was 2328.2±0.90.7 MPa and the hardness value was 81.3±0.1 HRA, respectively. In particular, when the optimally sintered Vanadis 60 composite underwent sub-zero plus heat treatments, the TRS and hardness values obviously increased to 2456.6±76.3 MPa and 85.7±0.3 HRA, respectively. Finally, the TEM and EBSD study results also revealed that various MC and M6C-type carbides significantly appeared in the Vanadis 60 composite after vacuum sintering and sub-zero plus heat treatments.
High-speed steels are materials used in cutting tools for machining other materials. The traditional high-speed steels contain a high amount of carbon and other alloying elements with different functions.1,2) Notably, Uddeholm Vanadis 60 is a high-alloyed, high-performance powder metallurgy (PM) high-speed steel with an addition of cobalt (10.5%). It is particularly suitable for cold work tooling where the highest wear resistance and compressive strength are required simultaneously.3,4) Metal matrix composites (MMCs) are high-strength and high-hardness materials with superior mechanical properties. Different matrixes were studied, such as aluminum, stainless steel, and high-speed steel matrix composites, which showed outstanding wear behavior. Furthermore, the use of PM techniques for manufacturing high-speed steel composites is increasing, which includes mixing the original powders, followed by sintering and consolidation.5,6,7)
In addition, these composites typically use ceramic phases as reinforcing phases, such as TaC, NbC, WC, B4C, or carbon fibers, which are dispersed in a matrix phase and produce composites with exceptional mechanical strength and toughness. These carbides are widely used as grain growth inhibitors in iron-based alloys.8,9,10,11) Among them, tantalum niobium carbides (TaNbC) combine the advantages of both tantalum carbides (TaC) and niobium carbides (NbC), which contribute significantly to improving the strength and hardness of steel composites.12,13)
Additionally, boron carbide (B4C) can effectively increase the relative density of steel composites. This is owing to a positive effect of B4C, which resembles a benignant effect taken by an “in situ” synthesis on metal matrix composites.14,15,16) According to the above literatures report, the TaNbC–B4C carbides are good choices to reinforce iron-based alloys. Thus, this study utilized the Ta50Nb50C–B4C carbides to strengthen the mechanical properties of Vanadis 60 high-speed steels.
Generally, sub-zero treated materials contain significantly reduced retained austenite (γR) amounts as a result of the isothermal and time-dependent martensite transformation, which takes place during the sub-zero treatment, and is an effective way to improve and enhance the properties of high-speed steels.17,18,19) Thus, sub-zero treatment shall help improve the properties of Vanadis 60 composite.
In our prior study, Vanadis 60 high-speed steel powders incorporated with 1%Ta50Nb50C provided high sintered density and improved mechanical properties after sintering.13,20) To further enhance the properties, we focused on addition of Ta50Nb50C–B4C to Vanadis 60. Hence, the aim of this study was to evaluate a series of vacuum sintering processes, and sub-zero and heat treatments for Vanadis 60 high-speed steel composites and further investigate the effects on the sintering behavior, strengthening mechanism, and mechanical properties of the Ta50Nb50C–B4C carbide-strengthened Vanadis 60 high-speed steel composite.
In this study, Vanadis 60 high-speed steel powders were employed as the substrate, and varying proportions of Ta50Nb50C–B4C powders were added as the strengthening phase. The Vanadis 60 composite was sintered from 1205 to 1225°C (more than 1230°C will produce a melting phenomenon). Meanwhile, the chemical compositions (mass%) of the Vanadis 60 high-speed steel powders were, as follows: 2.3% of C, 4.2% of Cr, 7.0% of Mo, 6.5% of W, 6.5% of V, 10.5% of Co, and 63.0% of Fe.
In order to further explore the effects of adding different Ta50Nb50C–B4C carbides to Vanadis 60, various ratios of Ta50Nb50C–B4C powders were added, and the addition of (X-Y)Ta50Nb50C–YB4C was utilized in this paper. Where, X equals 1 (as mentioned earlier, the optimal added amount of Ta50Nb50C powders was 1 mass%), and Y values were 0.025, 0.05, 0.1, and 0.2, respectively. In other words, the added Ta50Nb50C amounts were 0.975, 0.95, 0.9, and 0.8 mass%, and the amounts of B4C were 0.025, 0.05, 0.1, and 0.2 mass%, respectively. Thus, these composite specimens are hereafter designated as 0.025B (0.975 mass% Ta50Nb50C-0.025 mass% B4C), 0.05B (0.95 mass% Ta50Nb50C-0.05 mass% B4C), 0.1B (0.9 mass% Ta50Nb50C-0.1 mass% B4C), and 0.2B (0.8 mass% Ta50Nb50C-0.2 mass% B4C).
In this paper, the mean particle size of the Vanadis 60 high-speed steel powders was 28.1±0.5 μm, and the powders possessed a uniform and spherical appearance, while the mean particle sizes of the Ta50Nb50C and B4C powders were about 2.9±0.2 and 0.7±0.1 μm, respectively. The shape of the Ta50Nb50C and B4C powders was an irregular polygon. The mixed Vanadis 60 composite powders were milled using WC balls for 6 hours. The green compact (40×6×6 mm3) of the powder specimen was produced under uniaxial pressure of 300 MPa for 300 s.
To evaluate the effects of a series of vacuum sintering, sub-zero and heat treatments, the vacuum sintering temperatures were set at 1205, 1215, and 1225°C for 1 h in a 5×10−3 Pa, respectively. At the same time, heat treatment (quenching followed by tempering, designated as HT) and sub-zero plus heat treatment (quenching followed by sub-zero and tempering, designated as SZ) were performed, in which the samples were heated to 1180°C and the temperature was maintained for 30 min for quenching with 0.8 MPa of N2 as the quenching medium. The sample should be sub-zero treated (a temperature of −150°C for 50 min) followed by tempering.3) The tempering temperature was held at 550°C for 3 h, then cooled to room temperature, and repeated three times.
Furthermore, various material characterization techniques were used to evaluate the Vanadis 60-Ta50Nb50C-B4C high-speed steel matrix composite, which included apparent porosity (following the ASTM C830 standard) and mean grain size. Moreover, X-ray (D2 PHASER) analysis was used to examine the retained austenite of Vanadis 60 composites (following the SAE SP-453 and ASTM E975).7,9,13) Microstructural observations of the specimens were performed by optical microscopy (OM, Nikon LV-150N), scanning electron microscopy (SEM, Hitachi-S4700), transmission electron microscopy (TEM, JEOL JEM-2100F) and electron backscatter diffraction (EBSD, JSM-7800F).
Hardness tests were performed by Rockwell A hardness (HRA, Indentec 8150LK) with a loading of 588.4 N, which followed the ASTM B294 standard. While the Hung Ta universal material test machine (HT-2402) with a maximum load of 245 kN was used for the transverse rupture strength (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, and F is the maximum fracture load. Where the L was 30 mm, k was the 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 each data point was measured with at least three samples to get the average value.
Figure 1 reveals the apparent porosity, volume shrinkage, and mean grain size at different sintering temperatures for various ratios of Ta50Nb50C–B4C added to the Vanadis 60 specimens. The experimental results show that after adding B4C, the apparent porosity of the specimen was greatly reduced, and the measurement values were all less than 0.3%, as shown in Fig. 1(a). This confirms that the addition of B4C indeed promotes the formation of the liquid phase during the sintering process, thereby reducing the apparent porosity of Vanadis 60 composite. Among them, the 1225°C-sintered 0.025B specimen possessed the lowest apparent porosity (0.11±0.04%).
The volume shrinkage is also shown in Fig. 1(a). It can be observed that the 1215°C-sintered 0.025B specimen, with a small amount of B4C added and a low sintering temperature, achieved the highest volume shrinkage value (36.58±0.52%). It is worth noting that, at the same sintering temperature, the volume shrinkage rate showed a downward trend with the increased B4C addition. These results indicate that adding a small amount of B4C effectively increased the content of the liquid phase during the sintering process, which is consistent with previous literature.21) Interestingly, when the amount of B4C was increased, the content of the liquid phase also increased, resulting in reduced volume shrinkage of the specimens.
Based on the above two analysis results, it can be known that adding a small amount of B4C effectively increases the volume shrinkage and decreases the apparent porosity, which potentially leads to the Vanadis 60 composite having improved mechanical properties. However, it is important to note that the apparent porosity of the specimen was already very low under all sintering parameters. Consequently, porosity changes may have limited impact on the mechanical properties of this material; therefore, it remains necessary to conduct subsequent measurements and further analysis to determine the optimal sintering parameters.
The mean grain size of the Vanadis 60-Ta50Nb50C-B4C samples at various proportions and different sintering temperatures was determined using the linear intercept method,3,22,23) and the results are shown in Fig. 1(b). Notably, the 0.025B specimen sintered at 1205°C exhibited the smallest grain size, measuring 7.96±0.73 μm. As the sintering temperature and B4C additions increased, the formation of a liquid phase during the sintering process became more favorable, leading to an increase in the mean grain size. Among these observations, even when sintered at the lowest temperature of 1205°C, the mean grain sizes of the 0.1B and 0.2B specimens were significantly larger, as compared to the 0.025B specimen sintered at the same temperature. However, the mean grain sizes of the 0.025B and 0.05B specimens exhibited only a slight grain growth phenomenon until 1225°C, which suggests that all the specimens entered the stage of grain coarsening during sintering at 1225°C. In essence, further elevating the sintering temperature might surpass the optimum range and result in excessive grain growth.
Figure 2 illustrates the hardness and TRS test results for the 0.025B, 0.5B, 0.1B, and 0.2B specimens after different sintering temperatures. The experimental values for hardness and TRS are listed in Tables 1 and 2, respectively. Based on the former mean grain size test results, it is evident that the 0.025B and 0.5B specimens sintered at 1205°C exhibited relatively small grain sizes (7.96±0.73 and 11.23±0.74 μm, respectively). Additionally, the Hall-Petch equation verified that as the grain size decreased significantly, hardness tended to increase;24,25,26) therefore, the 1205°C-sintered 0.025B and 0.5B samples exhibited higher hardness values of 81.3±0.1 and 81.4±0.1 HRA, respectively.
| Temp. (°C) | Vanadis 60 | 0.025B | 0.5B | 0.1B | 0.2B | Vanadis 60+HT | Vanadis 60+SZ | 0.025B+HT | 0.025B+SZ |
|---|---|---|---|---|---|---|---|---|---|
| 1205 | − | 81.3±0.1 | 81.4±0.1 | 79.1±0.2 | 78.8±0.1 | − | − | 85.6±0.2 | 85.7±0.3 |
| 1215 | − | 81.5±0.1 | 81.6±0.1 | 78.6±0.3 | 78.4±0.1 | − | − | − | − |
| 1220 | 77.8±0.4 | − | − | − | − | 85.0±0.3 | 85.1±0.1 | − | − |
| 1225 | − | 80.8±0.1 | 79.9±0.2 | 78.2±0.4 | 78.1±0.3 | − | − | − | − |
| 1235 | 78.6±0.2 | − | − | − | − | − | − | − | − |
| 1250 | 78.6±0.2 | − | − | − | − | − | − | − | − |
| Temp. (°C) | Vanadis 60 | 0.025B | 0.5B | 0.1B | 0.2B | Vanadis 60+HT | Vanadis 60+SZ | 0.025B+HT | 0.025B+SZ |
|---|---|---|---|---|---|---|---|---|---|
| 1205 | − | 2328.2±90.7 | 2293.5±122.9 | 1413.6±37.3 | 1335.6±76.7 | − | − | 2241.5±84.3 | 2456.6±76.3 |
| 1215 | − | 2251.4±73.6 | 2217.7±51.7 | 1349.5±35.7 | 1018.9±62.8 | − | − | − | − |
| 1220 | 1842.3±100.8 | 2204.7±24.5 | 2185.3±30.4 | 1249.3±26.8 | 949.9±25.8 | 1317.8±119.0 | 1447.9±157.2 | − | − |
| 1225 | − | − | − | − | − | − | − | − | − |
| 1235 | 1685.9±45.2 | − | − | − | − | − | − | − | − |
| 1250 | 1484.7±78.7 | − | − | − | − | − | − | − | − |
The highest hardness, with a value of 81.6±0.1 HRA, was observed in the 1215°C-sintered 0.5B specimens, as depicted in Fig. 2(a). It was presumed that this sample did not undergo grain-coarsening, and thus, maintained a mean grain size of approximately 13.57±1.56 μm. Throughout this study, it was observed that as the sintering temperature (up to 1215°C) or B4C addition (0.1B or 0.2B) continued to increase, carbides gradually accumulated at the grain boundaries and underwent growth, thus, the strengthening effect of the carbides on the matrix phase was gradually diminished. As a result, the hardness of Vanadis 60-Ta50Nb50C-B4C specimens exhibited a declining trend.
Figure 2(b) shows the TRS test results for the 0.025B, 0.5B, 0.1B, and 0.2B specimens after different sintering temperatures. Notably, an increase in sintering temperature led to a significant decrease in all TRS values. Given that the apparent porosity of all the specimens was less than 1%, the differences were relatively small. The main factor affecting the TRS value of the Vanadis 60 composite was the mean grain size, as well as the distribution and morphology of the carbides (which will be further confirmed through subsequent OM and SEM observations). The highest TRS value was observed in the 1205°C-sintered 0.025B sample, which reached 2328.2±90.7 MPa. As mentioned earlier, this specimen possessed low apparent porosity (0.21±0.03%) and the smallest grain size (7.96±0.73 μm), which contributed to its elevated TRS value. Upon adding more B4C or increasing the sintering temperature, the content of the liquid phase during sintering increased;21) consequently, the effect of fine grain strengthening was diminished, which resulted in a downward trend in TRS, and this trend aligned with the results observed in the hardness tests.
Figure 3 displays the OM images of the 1220°C-sintered Vanadis 60 and the 1205°C-sintered 0.025B, 0.5B, 0.1B, and 0.2B specimens, respectively. As seen in Figs. 3(a) and 3(b), it is evident that two main types of carbides were precipitated within the Vanadis 60 substrate and Vanadis 60 composite specimens: Firstly, there are gray spherical or elliptical carbides (MC-type) distributed along the grain boundaries and within grains. Secondly, there are white plate-like carbides (M6C-type) precipitated both along the grain boundaries and in lumps within the grains. Based on previous literature,13) it is reasonable to hypothesize that these two types of carbides correspond to MC and M6C-type carbides, respectively. Specifically, MC-type carbides are mainly precipitated by vanadium (V); while M6C-type carbides are associated with precipitation by iron (Fe), tungsten (W), and molybdenum (Mo). Subsequent analysis involving EDS, EBSD, and TEM techniques will be conducted for a more in-depth discussion and confirmation.
As shown in Figs. 3(b), 3(c), 3(d), and 3(e), the internal pores within the four samples were nearly eradicated, which is indicative of successful densification at this sintering temperature. Notably, even at the lowest sintering temperature of 1205°C, signs of grain coarsening during the sintering process were already evident in the 0.1B and 0.2B specimens, as shown in Figs. 3(d) and 3(e). Concurrently, these carbides precipitated along the grain boundaries, which significantly compromised the material’s mechanical properties. This clearly suggests an excessive amount of B4C was added to Vanadis 60 composite at this temperature. Conversely, the 0.025B and 0.05B specimens exhibited finer grains, as shown in Figs. 3(b) and 3(c). It is possible to say that the addition of boron increased the liquid phase content during the sintering process, which effectively enhanced the sintering density, and this enhancement contributed to the improved mechanical properties.
Based on the above results and discussion, it is evident that the 1205°C-sintered 0.025B specimens exhibited both the highest TRS (2328.2±90.7 MPa) and elevated hardness (81.3±0.1 HRA). Therefore, this study selected these optimally sintered specimens for subsequent HT and SZ processes. Figure 4 displays the SEM images of the 1205°C-sintered 0.025B, 0.025B+HT, and 0.025B+SZ specimens. Earlier literature indicated that boron can increase diffusion rates by reducing the activation energy required for diffusion, particularly for atoms along grain boundaries.27) Therefore, the addition of B4C can provide activation energy greater than that needed for the LSP between grains, even at lower temperatures, and this effect accelerates the completion of sintering densification. As previously mentioned, two primary types of carbides precipitate within the Vanadis 60 composites; the first type consists of MC-type carbides, which are predominantly composed of V elements and distributed within the matrix; the second type appears as a white, thin stripe that precipitates along the grain boundaries, and represents M6C carbides mainly composed of Fe, W, and Mo elements. As indicated by the arrows in Fig. 4(a), these carbide precipitations are capable of contributing to the subsequent enhancement of mechanical properties.
Figure 4(b) reveals the SEM image of the 1205°C-sintered 0.025B specimens after heat treatment, where it is evident that the microstructure still consists of tempered martensite, MC-type, and M6C-type carbides within the matrix phase. Additionally, finer carbide precipitation was observed after the sub-zero treatment, and these refined carbides were more evenly distributed, as shown in Fig. 4(c). Further EDS analysis confirmed that they were MC-type carbides rich in V, and M6C-type carbides rich in W, Mo, and Fe elements, respectively. Among them, the irregularly shaped and refined carbides (Location 1, as seen in Fig. 4(c)) distributed in the grains and matrix phase had a high content of V (27.03 at% of V) carbide, as listed in Table 3(c)-1. It was reasonable to surmise that these carbides predominantly belong to MC-type carbides (rich in V). Furthermore, the white thin stripe carbides (Location 2, as seen in Fig. 4(c)) within the grains and grain boundaries contained high contents of W-rich, Mo-rich, and Fe-rich carbides (11.10 at% of W, 17.42 at% of Mo and 22.87 at% of Fe), as listed in Table 3(c)-2. It was reasonable to surmise that these carbides mainly belong to M6C-type carbides (rich in W, Mo, and Fe). Nevertheless, this paper will conduct further confirmation of these two primary precipitates using EBSD and TEM techniques.
| Elements (at%) | (c)-1 | (c)-2 |
|---|---|---|
| C | 51.88 | 35.25 |
| V | 27.03 | 3.32 |
| Cr | 2.90 | 4.10 |
| Fe | 1.75 | 22.87 |
| Co | 0.47 | 4.17 |
| Nb | 1.21 | 0.81 |
| Mo | 9.56 | 17.42 |
| Ta | 0.42 | 0.96 |
| W | 4.78 | 11.10 |
Figure 5(a) displays the XRD patterns of the 1220°C-sintered Vanadis 60, Vanadis 60+HT, and Vanadis 60+SZ specimens, as well as the 1205°C-sintered 0.025B, 0.025B+HT, and 0.025B+SZ specimens. All Vanadis 60 and 0.025B specimens exhibited evident phase transformations and carbide precipitation after sintering, heat treatment (HT), and sub-zero treatment (SZ). These carbides predominantly consist of MC-type carbides precipitated by V, and M6C-type carbides precipitated by Mo, W, and Fe. Furthermore, the intensity values of the austenite peak for the Vanadis 60+SZ and 0.025B+SZ samples were relatively low, indicating that the austenite phase of these specimens could disappear, and be mostly replaced by a complete martensite phase. The retained austenite content of the Vanadis 60 and 0.025B-sintered specimens significantly decreased after sub-zero treatment, from 5.12±1.64 to 0.78±0.37 vol.%; and from 3.83±0.35 to 0.560±0.12 vol.%, respectively. Consequently, the XRD analysis confirmed that sub-zero treatment effectively reduced the retained austenite, which contributed to the enhanced mechanical properties of the Vanadis 60-Ta0.5Nb0.5C-B4C composites.
Figure 5(a) displays the XRD patterns utilized to calculate the retained austenite content in Vanadis 60 and 0.025B-sintered, heat treated, and sub-zero treated samples, and the resulted volume fraction of retained austenite is displayed in Fig. 5(b). The 1220°C-sintered Vanadis 60 specimen exhibited a retained austenite content of 5.12±1.64 vol.%, while the 1205°C-sintered 0.025B specimen had 3.83±0.35 vol.% retained austenite. The high cobalt content (10.5%) in Vanadis 60 effectively facilitates the transformation of austenite into martensite within iron-based composites.13) Additionally, the addition of B4C consumed the alloy content in the matrix phase, which led to increased carbide formation, thereby reducing the retained austenite content, and promoting martensite transformation.27) Our previous studies indicated that retained austenite can be eliminated through sub-zero (−150°C) and high-temperature tempering (550°C) treatments.7,9,13) Thus, the retained austenite contents in the Vanadis 60+SZ and 0.025B+SZ samples were significantly decreased to 0.78±0.37 and 0.56±0.12 vol.%, respectively, which reinforces the effectiveness of sub-zero treatment in decreasing retained austenite content within the Vanadis 60 composite.
Figure 6 presents a comprehensive comparison of the mechanical properties of the different treatments of the Vanadis 60 and Vanadis 60 composites, and Tables 1 and 2 show the complete data regarding hardness and TRS. Obviously, the hardness values of Vanadis 60+HT and Vanadis 60+SZ are significantly higher than that of the sintered Vanadis 60 specimens (without heat treatment). Among them, Vanadis 60+SZ demonstrates the highest hardness value, reaching 85.1±0.1 HRA, as shown in Fig. 6(a). Remarkably, the 0.025B specimens exhibited a substantial increase in hardness after heat and sub-zero treatments, measuring 85.6±0.2 HRA and 85.7±0.3 HRA, respectively. This phenomenon is likely attributed to quenching and tempering during the heat treatment, which led to the solid dissolution of primary carbides back into the matrix phase, followed by their re-precipitation. This process causes deformation in the crystal lattice and results in the secondary hardening effect. Additionally, the secondary carbides, such as MC and M6C, possessed high hardness and uniform distribution, which significantly enhanced the hardness of the specimens.
Comparisons of the TRS values for different treatments of Vanadis 60 and Vanadis 60 composite specimens are shown in Fig. 6(b), with detailed data presented in Table 2. Notably, an upward trend in TRS values was observed for the 0.025B specimens after undergoing sub-zero plus heat treatment (SZ). The initial TRS value for the 0.025B specimens was 2328.2±90.7 MPa, with the highest TRS value achieved by the 0.025B+SZ specimens (2456.6±76.3 MPa). In contrast, the TRS value of the Vanadis 60+HT specimen exhibited a distinct downward trend (1842.3±100.8 → 1317.8±119.0 MPa) following quenching and the tempering heat treatment.
Although a minor TRS increase was observed after the sub-zero treatment (1447.9±157.2 MPa), it remained lower than the TRS of the 1220°C-sintered Vanadis 60 specimens. This trend can be attributed to the secondary hardening experienced by Vanadis 60 high-speed steel during high-temperature tempering, which led to hardness and brittleness. Moreover, the substantial alloy and carbon content in Vanadis 60 high-speed steel contributed to the increased hardness post-heat treatment, accompanied by enhanced carbide precipitation, which in turn elevated the brittleness of the material.28) Notably, these findings align with our prior research.13)
Figure 7 shows the carbide volume percentages (V%) resulting from the different heat treatment processes. The 1220°C-sintered Vanadis 60 specimen comprised a 79.7% matrix phase and 20.3% carbides. Among the carbides, the main MC-type carbides account for 17.1%, while the M6C-type carbides account for 3.2%. The addition of Ta50Nb50C-B4C powders introduced Nb and B elements that facilitated the formation of M6C-type carbides, which led to an increased content of M6C-type carbides. As a result, the 1205°C-sintered 0.025B specimen comprised a 76.5% matrix phase, 16.7% MC-type, and 6.8% M6C-type carbides.
More alloying elements dissolved and re-precipitated into the matrix phase through the HT and SZ processes, which caused a reduction in MC and M6C-type carbide content. The 0.025B+HT specimen consisted of a 77.8% matrix phase, 12.1% MC-type, and 10.1% M6C-type carbides. Similarly, the 0.025B+SZ specimen included an 82.2% matrix phase, 10.2% MC-type, and 7.6% M6C-type carbides. Overall, the M6C-type carbide content in the 0.025B+SZ specimen (7.6%) significantly surpassed that in the sintered Vanadis 60 specimen (3.2%). Generally, MC-type carbides are mainly spherical or elliptical shapes composed of vanadium. In contrast, M6C-type carbides are generally plate-like shapes (rich in molybdenum, tungsten, and iron, as shown in Fig. 3(b)). When high-speed steel is tempered, the carbides precipitated in the form of MC-type can significantly improve the wear resistance and thermal stability of high-speed steel. However, carbides precipitated in the form of M6C-type (normally are (Fe, W)6C carbides precipitated during tempering) are the most critical factor in the secondary hardening of high-speed steel. Relatively speaking, M6C-type carbides can more effectively enhance the anti-tempering stability of high-speed steel. In other words, M6C-type carbides can more effectively increase hardness and strength than MC-type precipitates.7,13,29,30,31) This finding suggests a positive relationship between the increased hardness and TRS and the quantity of M6C-type carbides.
As previously analyzed and discussed, the sub-zero plus heat treatment (SZ) proved highly effective in reducing the retained austenite contents. Furthermore, the re-precipitation of refined MC-type carbides within the grains led to a marked enhancement in TRS. Consequently, the incorporation of an appropriate amount of Ta50Nb50C-B4C powders (as observed in the 0.025B+SZ specimen) combined with the application of sub-zero plus heat treatment collectively contributed to a notable improvement in the material strength of Vanadis 60 high-speed steel.
Figure 8 reveals the EBSD observations of the 1205°C-sintered 0.025B+SZ specimens. Notably, the primary carbide precipitates in these specimens are MC and M6C-type carbides, and among these, the most abundant are the refined MC-type carbides (white precipitates) that re-precipitated within the grains. Irregular M6C-type carbides were observed to precipitate along the grain boundaries, as shown in Figs. 8(a) and 8(b). It is of particular interest to observe that drastic changes in quenching temperature or phase transformations can induce significant internal stress, and this phenomenon prompted the extrusion of MC-type carbides from the two ends toward the central M6C-type carbides. This unique process resulted in the formation of a distinct structure wherein the three kinds of carbides interconnected.
Figure 8(c) displays the grain orientation analysis of the 1205°C-sintered 0.025B+SZ specimens, and indicates that the carbides exhibited no special crystallographic orientation. This observation is consistent with the characteristic behavior of PM materials, which typically lack orientation. Based on the EBSD analysis results (Fig. 8(d)), it is evident that the carbides precipitated in the Vanadis 60 composite were mainly comprised of MC and M6C-type carbides. Furthermore, Fig. 8(e) represents the pole figure of EBSD analysis for the 1205°C-sintered 0.025B+SZ specimens. The examination of the crystal orientation in the iron-based substrates shows a relatively uneven distribution of hot spots, indicating the absence of any specific directionality during the sintering process. Notably, the pole figure of Fe-BCC displayed hot spots on the {100}, {110}, and {111} planes (Fig. 8(e)), further validating the absence of pronounced crystallographic orientation within the Vanadis 60 composite. This experimental result supports the argument that PM-produced materials exhibit no inherent orientation.
Based on the preceding discussion and results, the primary precipitates in the 1205°C-sintered 0.025B+SZ specimens were identified as MC and M6C-type carbides. To gain deeper insights into their microstructure, TEM analysis was employed, as shown in Figs. 9 and 10. Within this figure, Figs. 9(a), 9(b), and 9(c) correspond to the selected area electron diffraction (SAED), bright field, and dark field images of the MC-type carbides, respectively.
The SAED image in Fig. 9(a) reveals diffraction patterns that are characteristic of a cubic face-centered cubic (FCC) crystal structure, specifically diffracted along the [011] crystal axis. This finding confirms the MC-type carbide’s crystalline arrangement is a sodium chloride (NaCl) structure within the Fm3m cubic crystal system and the associated space group. Additionally, the morphological analysis of the MC-type carbide was irregular, as observed in Figs. 9(b) and 9(c), which indicates an irregular shape that is consistent with our earlier discussion in this paper.
According to Bragg’s law and the following equation,32)
Where, d is the interplanar spacing; a is the lattice parameter; and h, k, and l are the Miller indices. The interplanar spacing of the (200) plane of VC can be theoretically calculated as 0.208 nm. However, upon examining Fig. 9, it became evident that different alloying elements were incorporated into the MC-type carbides, which caused an expansion of the interplanar spacing. Consequently, the actual interplanar spacing of the MC-type carbides (200) in this experimental setup was measured to be 0.222 nm, and this slight variance between the experimental and calculated values can be attributed to these factors. The TEM observations offer further verification that the MC-type carbide formation was a result of the solid-solution involving the V element.
Figure 10 shows the TEM observations of the M6C carbides found in the 1205°C-sintered 0.025B+SZ specimens. As evidenced by Figs. 10(b) and 10(c), the M6C-type carbide exhibited an irregular and elliptical morphology. Calculations based on the SAED image in Fig. 10(a) indicate that it corresponds to a cubic FCC structure, diffracting in the direction of the [001] crystal axis. This verifies that the M6C-type carbide possessed a diamond structure characteristic of the Fd3m cubic crystal system within the designated space group. Considering that Mo, W, and Fe elements give rise to carbides with the same structural attributes. it can be computed that the interplanar spacing for the (220) plane of Fe3W3C (or Fe3Mo3C) equates to 0.391 nm; however, the (220) interplanar spacing of the M6C-type carbide in this particular experiment was observed to be 0.305 nm. This discrepancy can be attributed to the presence of numerous other alloying elements within M6C-type carbides. As it is easy to cause the interdiffusion of different alloy atoms in the carbide, the experimentally measured interplanar spacing changed greatly. Nevertheless, it was reasonable to conclude that the M6C-type carbide was formed by the solid solution of W, Mo, and Fe elements.
In this research, TEM observations further confirmed the presence of MC-type carbides, which precipitated with a sodium chloride (NaCl) structure of the Fm3m cubic crystal system in the space group. These precipitates were formed due to the solid solution of the V element. Additionally, M6C-type carbides were observed, which precipitated with the diamond structure of the Fd3m cubic crystal system in the space group. The interplanar distance in these M6C-type carbides was identified as Fe3W3C (or Fe3Mo3C), as resulted from the solid solution of W, Mo, and Fe elements. It should be noted that other alloying elements may have also been present in the MC and M6C-type carbides, which led to the slight discrepancies between the experimental and calculated values of interplanar spacing in this paper. However, despite the potential presence of other alloying elements, overall, the TEM observations firmly established the existence of both MC and M6C carbides in the Vanadis 60-Ta0.5Nb0.5C-B4C high-speed steel composite after undergoing vacuum sintering, sub-zero, and heat treatments. These observations contributed to a better understanding of the microstructure and carbide formations in the composite material, and aided in the evaluation of its properties and potential applications.
The results of this study reveal that excellent mechanical properties were achieved through the addition of 0.975 wt% Ta50Nb50C-0.025 wt% B4C powder (0.025B specimen) sintered at 1205°C for 1 hour. The apparent porosity was measured at 0.21±0.03%, the TRS value was 2328.2±90.7 MPa, and the HRA hardness value reached 81.3±0.1 HRA. Furthermore, subsequent heat treatment (0.025B+HT) and sub-zero treatments (0.025B+SZ) led to significant improvements in TRS values, which increased to 2241.5±84.3 and 2456.6±76.3 MPa, respectively. These treatments also yielded higher hardness values, measuring 85.6±0.2 and 85.7±0.3 HRA.
The sub-zero treatment effectively reduced the retained austenite content (from 3.83±0.35 to 0.56±0.12 vol.%), thereby further enhancing the TRS of Vanadis 60-Ta0.5Nb0.5C-B4C composites (0.025B+SZ). Moreover, the 0.025B+SZ specimen consists of 82.2% matrix phase, 10.2% MC-type, and 7.6% M6C-type carbides. Notably, the M6C-type carbide content significantly exceeded that of the sintered Vanadis 60 specimen (3.2%); hence, it was reasonable to conclude that the observed increased hardness and TRS values exhibited a positive relationship with both the reduced retained austenite content and the higher proportion of M6C-type carbides.
The TEM and EBSD research results further confirmed the presence of MC and M6C-type carbides in the specimens after vacuum sintering, heat treatment, and sub-zero treatment. The MC-type carbides were found to be rich in V elements, whereas the M6C-type carbides were predominantly composed of W, Mo, and Fe elements. The TEM observations also verified that the MC and M6C-type carbides exhibited a NaCl and diamond structure, respectively.
This research is supported by the National Science and Technology Council of the Republic of China under Grant No. NSTC 112-2221-E-027-035-. The authors would like to express their appreciation for voestalpine High Performance Metals Pacific Pte Ltd. and ASSAB Steels Taiwan Co., Ltd. Furthermore, thanks to Prof. H.C. Lin and Mr. C.Y. Kao of Instrumentation Center, National Taiwan University for EBSD experiments. Moreover, thanks to Miss Chuang of Precision Analysis and Material Research Center, National Taipei University of Technology for TEM experiments.
