Fabricating and Characterization of MPEA Binder Phase Cemented Carbide and Its Comparison with WC-Co

Abstract

The development and research of physically superior multi-principal element alloy (MPEA) binders as cemented carbide binders is a hot topic. In this work, we fabricated a new type of MPEA binder-cemented carbide using the powder metallurgy method and investigated the effects of ball milling parameters and sintering temperature on the microstructure and mechanical properties of the cemented carbide. The results are compared with those of cobalt binder samples under the same conditions. The results show that the ball milling parameters for low-speed long ball milling time are superior to those for high-speed low ball milling time. Compared with the pure cobalt binder, MPEA binder-cemented carbide significantly slows down the growth of WC grains, improves the mechanical properties of cemented carbide, and achieves a combination of TRS of 2741.5 MPa and Rockwell hardness of 91.1 HRA. The multi-principal element alloy (MPEA) binder has the potential to become an excellent substitute for Co.

1. Introduction

WC-Co-cemented carbides are one of the most widespread powder metallurgy products composed of one or more refractory metal carbides as hard grain phases (WC, TiC, etc.) and soft ductile phases as binders (Fe, Co, Ni, etc.), which have long been used as cutting tools, dies, and wear-resistance parts owing to their elevated hardness, strength, elastic modulus, rigidity, and wear resistance [1,2,3,4]. However, with the rapid development of the industry, traditional WC-Co-cemented carbides are unable to meet the increasing demand for mechanical properties. Therefore, it is of great significance to prepare new cemented carbides with high hardness and strength.

In recent years, with the prosperity of the new energy industry, the demand for Co is significantly increasing. Cobalt is an important strategic resource element in China, the price of which is rapidly increasing [5,6]. When Co is used as a binder, increasing the mass fraction of the binder phase will lead to an increase in alloy toughness and a decrease in strength and hardness, which is the “irreconcilable contradiction” of traditional WC-Co-cemented carbides [7,8]. Research on using other alloys to replace the Co binder has promising prospects.

Many studies have been conducted, both domestically and internationally, on using MPEA as a binder of cemented carbide. The advantage of the MPEA binder over Co is that it can obtain more comprehensive physical properties and excellent high-temperature mechanical properties through composition design. Erik et al. [9] used Thermo-Calc Software and EMTO (Exact Muffin tin Orbitals) methods to perform thermodynamic calculations on the phase stability of (CoNiFeCr)_(1−x) WxHEA. This method combines the phase diagram with thermodynamic properties information, using Gibbs free energy as a function of temperature, pressure, and composition for all possible phases, seeking the lowest point of the entire system energy to calculate the relatively stable bonding phase. Yadav et al. [10] utilized the Integrated Computational Materials Engineering (ICME) method to design and develop a CoNiFeCrMo binder as an alternative to Co. Preliminary studies have shown that Co-Ni-Fe as a binder can enhance the fracture toughness of WC ceramics, and poor oxidation resistance is the main drawback of the pure Co binder. Adding Cr can effectively enhance the oxidation resistance of cemented carbides, while Mo is beneficial for suppressing grain growth and solid solution strengthening.

MPEAs have a “ sluggish diffusion” effect, which can affect the interaction between WC grain and the binders to some extent, such as suppressing the diffusion of WC in the binders, thereby inhibiting the growth of WC grains, and suppressing other unnecessary carbides. The “cocktail” effect of MPEAs can be briefly described as the interaction between different elements causing the alloy to exhibit a composite effect. It emphasizes the role of the main elements of the alloy at the atomic scale, ultimately affecting the macroscopic properties of the alloy. The diversity of types, quantities, and element contents of high-entropy alloy components provides more choices for the composition design of cemented carbides binder alloys, which means that using MPEAs as binders can expand the use of cemented carbides [11,12,13]. The behavior of the binder phase during the pressure-sintering process will significantly affect the properties of cemented carbides. Therefore, the MPEA binder phase should have the following properties during the sintering process: the binder phase should not react with WC to form other carbides; the binder phase should have good wettability with WC grains; and the binder phase can reduce the sintering temperature and suppress grain growth [14,15]. Many studies have shown that the phase composition and bending strength of cemented carbides are sensitive to the carbon content. Excessive or insufficient carbon content greatly reduces the toughness or hardness of cemented carbides [16,17,18], which can also be learned from phase diagrams [19]. The liquid phase line of the binder phase will change with the change in composition. When the Ni component is increased, the liquidus line rises and the two-phase region widens. When the Fe content is increased, the liquidus line decreases and the width of the two-phase region narrows [20]. The two-phase regions of MPEA binders with different compositions will undergo varying degrees of left shift, which determines the microstructure and performance changes in MPEA binder-cemented carbides [21].

This paper uses WC powder with a carbon content of 6.13% as the hard phase to fabricate WC-CoNiFeCrCu and traditional WC = Co cemented carbides. The effects of different ball milling conditions and pressure-sintering temperatures on the microstructure, WC particle size, and mechanical properties of MPEA binder-cemented carbides were studied. Pure cobalt binder-cemented carbides were prepared using the same process, and the microstructure and mechanical properties were analyzed and compared.

2. Experimental

2.1. Materials and Processing

The WC powder and Co powder were provided by Zhuzhou Jwe Cemented Carbides Co., Ltd. (Zhuzhou, China), and the other metal powders were provided by Changsha TIJO Metal Materials Co., Ltd. (Changsha, China). The specific powder particle size is shown in

Table 1. Powder particle size.

Name	Size (μm)
WC	10
W	1.04
Co	1.27
Ni	1.96
Fe	2.3
Cr3C2	1

. The experiment set the ball milling time, ball milling speed, and sintering temperature changes. The specific milling parameter is shown in

Table 2. Ball milling parameters.

Sample	Milling Time (h)	Milling Speed (r/min)
A	10	200
B	15	200
C	20	200
D	45	105
E	50	105
F	55	105

. This article uses ethanol as the ball milling medium and adds the milling ball and a 2% mass fraction of PEG as a forming agent to the slurry. The ball-to-material ratio is 2:1, and the ball milling process grouping is shown in Table 2. The composition of MPEA binder cemeted carbide is shown in

Table 3. Composition of the MPEA binder-cemented carbides.

Powder	WC (wt%)	W (wt%)	Co (wt%)	Ni (wt%)	Fe (wt%)	Cr3C2 (wt%)
WC-MPEA	93.81	0.16	4.15	1	0.65	0.23

. After vacuum drying at 60 °C, the cemented carbides’ green bodies were pressed and formed and then liquid-phase-sintered in a gas pressure-sintering furnace for 100 min to obtain sintered samples. The sintering pressure was 5.8 MPa argon gas, and three sintering temperatures of 1410 °C, 1450 °C, and 1500 °C were set for each group of samples.

2.2. Microstructural and Physical Properties Characterization

Transverse rupture strength tests were conducted on cemented carbide samples according to the GB3851-2015 [22] standard using an electronic universal testing machine (INSTRON-5982 model (Instron, Boston, Massachusetts), with sample sizes of 6.5 mm × 5.25 mm × 20 mm. We used ImageJ software (ImageJ 1, National institution of Health, Bethesda, MD, USA) to calculate the WC grain size by counting 400–500 grains. SEM pictures were obtained with a Mira3 scanning electron microscope (TESCAN, Brno, The Czech Republic).

3. Results and Discussion

3.1. The Influence of Fabricating Process on the Microstructure and Properties of Alloys

To investigate the effect of ball milling parameters on the microstructure and properties of cemented carbides, groups ABC (milling time of 10, 15, and 20 h, respectively; 200 r/min) were set as the high-speed and low-speed ball milling time groups, and DEF (milling time of 45, 50, and 55 h, respectively; 105 r/min) was set as the low-speed and long ball milling time group. Figure 1 shows SEM images of three groups of samples under ball milling conditions ABC. It can be observed from Figure 1a,d,g that under the same sintering conditions of 1410 °C, the morphology of WC grains after 10–20 h of ball milling is circular, and the longer the ball milling time, the higher the degree of grain boundary fragmentation. Extending the ball milling time to 20 h under these sintering conditions can further increase the degree of grain fragmentation, but it does not reach the maximum grinding size at this speed. In Figure 1g,h,i, the grain body sintered at 1410 °C has a higher degree of fragmentation, but as the sintering temperature increases, the degree of WC grain fragmentation relatively decreases. When comparing Figure 1i with Figure 1c,f, the grains in Figure 1c,f are basically intact. This indicates that at a sintering temperature of 1500 °C, the dissolution and precipitation of WC grains can completely repair the grains after 10 and 15 h of ball milling. However, after 20 h of ball milling, the degree of fragmentation is too high and has not been repaired. The grains in Figure 1c,f are mostly trapezoidal and nearly triangular in shape. These results indicate that the dissolution and precipitation of WC in the binder phase simultaneously achieve the functions of grain growth and repairing grain boundaries [23,24,25].

From the data in

Table 4. Average grain size of some samples.

Group Name	Average Particle Size (μm)	Variance (σ2)
B 1410	5.933	1.932
B 1450	6.375	2.081
A 1500	9.299	2.434
B 1500	7.969	2.275
C 1500	7.371	2.162
E 1410	1.482	0.471
E 1450	1.390	0.661
D 1500	1.447	0.568
E 1500	1.465	0.465
F 1500	1.419	0.560

, it can be seen that as the sintering temperature increases, the grain size of Group B (milling time 15 h; 105 r/min) increases from 5.933, 6.375, and 7.969 μm, respectively, after the sintering temperature increases from 1410 °C to 1450 and 1500 °C. At 1500 °C, the ball milling time of groups A, B, and C was extended from 10 h to 15 h and 20 h, respectively, and the grain size decreased from 9.299 to 7.969 and 7.371 μm.

Based on the above experimental results, the ball milling parameters were adjusted to obtain better-performance cemented carbide samples. Three sets of ball milling conditions for DEF powders were set, with ball milling times of 45, 50, and 55 h, respectively, and ball milling speeds of 105 r/min. At the same sintering temperature of 1500 °C, the grain sizes of the DEF group were 1.4473, 1.4654, and 1.4193 μm, respectively. Compared with 9.299, 7.969, and 7.371 μm for the ABC groups at 1500 °C, the maximum reduction was about 6.55 times. From all SEM images of groups DEF in Figure 2, it can be seen that only the 1410 °C sintering temperature group of D (milling time 45 h; 105 r/min) displayed a relatively high degree of fragmented grain boundaries, as shown in the enlarged image of Figure 2a. Correspondingly, no fragmented grains were observed at the 1450 °C sintering temperature in group D, which increased the sintering temperature, and the 1410 °C temperature in group E, which increased the ball milling time. Correspondingly, no fragmented grains were observed at the 1450 °C sintering temperature in group D, which only increased the sintering temperature, and the 1410 °C temperature in group E (milling time of 50 h; 105 r/min), which only increased the ball milling time.

The line observed from the two-dimensional SEM image passing through the interior of the grain is actually an internal grain boundary in three-dimensional space. During the liquid-phase sintering process, the binder phase migrates and fills these internal grain boundaries, and its morphology is similar to the state when the grains are about to close in the later stage of sintering. There are two possibilities for fragmented grains to transform into intact grains:

• Through grain boundary diffusion, the binder in the pore surface diffuses toward the grain surface, causing the internal grain boundaries to disappear.
• The internal grain boundaries continue to extend, dividing a fragmented grain into two complete grains.

From Figure 3b, it can be observed that by extending the ball milling time, the surface energy of the WC grain system increases, resulting in an increase in the area of grain boundaries of the WC grains and the possibility that the cracks completely penetrate the original grains, dividing the grains into two parts and forming two complete grains. We observe the internal grain boundaries at the grain edge in Figure 3b, with sizes of 0.779 μm and 1.09 μm, respectively, which are close to the grain size of 1.48 μm. The microstructures in the left and right figures of Figure 3c are both intact grains, and the D (milling time 45 h; 105 r/min) 1410 °C sintering temperature group with similar ball milling conditions has cracks of about 1 μm, indicating that increasing the dissolution precipitation repair efficiency after sintering at 40 °C can repair the original cracks. After extending the ball milling time for 5 h, the grain size further decreased, making the original sintering temperature of 1410 °C sufficient to repair the cracks in the E (milling time 50 h; 105 r/min) 1410 °C sintering temperature group.

Figure 4 shows the TRS–Sintering Temperature plot and Rockwell hardness–Sintering Temperature plot of groups A–F of cemented carbide samples. The fracturing of cemented carbides is considered to be a two-step process. The Brittle inter- and trans-crystalline fracture of the carbide phase occurs first, followed by the ductile fracture of the bonded phase and the binder phase and WC–binder interface [26,27,28]. The area fraction of the crack path through the carbide phase accounts for at least 50–80% of the fracture surface area of cemented carbides and increases with the amount of carbide phase in the alloy [29]. The fracture resistance of cemented carbides is related to the free path of the binder phase and the distribution of WC grain boundaries. The continuous skeleton formed by fine WC grains increases the possibility of the crack tip meeting the grain boundary, often making transgranular fracture more favorable for crack propagation than intergranular fracture and binder fracture. It can be observed that the TRS of the DEF groups with intact grains and smaller grain sizes is much higher than that of the ABC groups with internal cracks and grain boundaries. In groups A–C, the TRS is less affected by the sintering temperature, while groups D–F are more affected by the temperature, and the TRS of group D (milling time 45 h;105 r/min) increases with the increase in the sintering temperature. This is because group D’s cemented carbide grains have more internal cracks and grain boundaries at the sintering temperature of 1410 °C, and the effect of repairing the cracks after raising the sintering temperature is more obvious. The TRS of group E (the 1450 °C sintering temperature group) was the highest, reaching 2741.5 MPa. The WC grains under this condition had a smaller grain size, better grain integrity, and higher TRS and hardness. Long-time low-speed ball milling compared with short-time high-speed ball milling can be maintained in the case of smaller grain edge porosity refinement of the grain so that the strength of the hardness is higher. Selecting the appropriate pressure-sintering temperature can repair the WC grain internal grain boundaries, and a sintering temperature of 1450 °C is enough to create the average WC internal grain boundary size, which is smaller than the E (milling time 50 h;105 r/min) group when completely closed.

3.2. Comparison of Microstructure and Physical Properties Between Cobalt and MPEA Binder Phase Cemented Carbides

Figure 5 shows the SEM images of the pure cobalt binder and MPEA binder-cemented carbide samples using Group E’s preparation conditions. Figure 5a–c shows the sintering temperature group of pure cobalt binder samples at 1410–1500 °C, and Figure 5d–f shows the sintering temperature group of MPEA binder-cemented carbides at 1410–1500 °C, all with a binder phase mass fraction of 6%. It can be observed that no η-phase is produced for either set of samples. The grain enlargement of the pure cobalt sample group with the increase in the sintering temperature is very obvious, and from 2.35 μm at 1410 °C to 3.641 μm at 1500 °C, the growth rate reaches 54%, and even larger grains above 20 μm appear. The size difference in the grains is large, the angle is frequent and obvious, and the size of the grains is not uniformly distributed. The MPEA binder sample groups are less affected, and the difference in grain size between the 1410 °C and 1500 °C groups is only about 1.1%, with no oversized grains. The rate of grain growth during sintering can be described by the following equation [30]:

$${\displaystyle \frac{dr}{dt}}={\displaystyle \frac{2D\gamma {\mathit{\Omega }}^2{C}_{\infty }}{RT{r}^2}}\times ({\displaystyle \frac{r}{\overline{r}}}-1)$$

(1)

In the formula, r is the grain size (nm), t is the sintering time, D is the solute atomic diffusion coefficient (m²/s), γ is the specific interfacial energy (J/m²), Ω is the molar volume (m³/mol), C_∞ is the average solubility of the solute (mol⁻¹), $\overline{r}$ is the average grain size at equilibrium (nm), T is the sintering temperature, and R is a constant. According to this equation, without considering other effects, the growth rate of grains is mainly related to the sintering temperature, holding time, solute solubility, and initial grain size. For MPEA binder-cemented carbides, some binder powders exhibit a liquid phase when melted at a lower sintering temperature of 1300 °C. Compared to pure cobalt samples, the binder in MPEA binder-cemented carbide samples contains Ni and Fe, and the solubility of WC is similar to that of pure cobalt binder samples. Therefore, the effect of solubility on the grain growth rate is relatively small. Under the same ball milling conditions, the initial grain size is the same, and other coefficients are also the same. Moreover, under sintering conditions, the grain growth degree of MPEA binder-cemented carbide samples is much smaller than that of the pure cobalt binder phase. At the sintering temperature of 1410 °C, as shown in Figure 5d, the original smooth-edge grains even still exist, indicating that the MPEA binder-cemented carbides are not only reduced in grain size due to the use of Cr as a grain inhibitor but also affected by the delayed diffusion effect of the multi-component alloy. The MPEA binder-cemented carbides undergo a large amount of mismatch at the microscale, meaning the atomic spacing is not constant and the diffusion-free path is reduced. This leads to a much slower diffusion rate of atoms dissolved in the MPEA binder during the sintering process of W and C elements compared to the pure cobalt binder, thereby inhibiting grain growth [31,32].

Table 5. Relative density of two different binders.

Group	Relative Density
Cobalt 1410	98.8%
Cobalt 1450	98.5%
Cobalt 1500	98.8%
MPEA 1410	97.1%
MPEA 1450	98.7%
MPEA 1500	98.4%

shows the relative density of samples of two different binders. It can be observed that the relative density of the samples is around 98%, which is a relatively high level.

To better observe and compare the distribution of elements during the sintering process, we selected samples at 1450 °C for EDS analysis. As shown in Figure 6c–g and

Table 6. EDS element total spectrum.

Element	Wt%	Wt% σ
O	0.57	0.08
Cr	0.28	0.04
Fe	0.82	0.05
Co	5.73	0.08
Ni	1.29	0.07
W	91.20	0.16

, the binding phase is composed of five elements: Co, Ni, Cr, Cu, and Fe, while there is no W in the enriched region of the bonding phase.

As shown in Figure 6e, although there is some aggregation of the bonding phases, it is not very severe, and with the increase in temperature, the atomic activation energy and diffusion ability increase, which is also beneficial for the distribution of bonding agents in hard alloys. From Figure 6, it can be seen that copper elements are uniformly distributed throughout the entire region, indicating the presence of some copper element precipitation.

According to Figure 7a, the TRS of the MPEA binder-cemented carbide sample reached 2741.5 MPa at a pressure-sintering temperature of 1450 °C, which is 388 MPa higher than the 2398 MPa of the pure cobalt binder-cemented carbide sample prepared under the same conditions, showing an increase of about 16%. The TRS of both binder groups first increased and then decreased with the increase in sintering temperature, so the optimal sintering temperature should be around 1450 °C. The hardness of pure cobalt binder-cemented carbides at a 1500 °C sintering temperature is only 88.78HRA, which is the result of the rapid expansion of grain size. According to

Table 7. Comparison of microscopic sizes of cemented carbides of two different binders.

Group	Grain Size (μm)
Cobalt 1410	2.350
Cobalt 1450	2.786
Cobalt 1500	3.641
MPEA 1410	1.482
MPEA 1450	1.390
MPEA 1500	1.465

, MPEA binder saples’ grain size significantly smaller than cobalt sample fabricated under same conditions.

4. Conclusions

(1)

The mechanical alloying method of low-speed long-term ball milling can reduce the edge pores of WC grains, which is beneficial to the integrity of WC grain structure and can effectively prevent the propagation of cracks in the alloy, thereby improving the TRS of the alloy.

(2)

Under the fabricating conditions of this experiment, the performance of the sample is relatively low at a sintering temperature of 1410 °C, while a sintering temperature of 1500 °C will cause excessive grain growth and reduce the mechanical properties of the hard alloy. A sintering temperature of 1450 °C can repair the grains while avoiding excessive growth, which is an equilibrium point.

(3)

Compared with WC-Co (2398 MPa, 90.4 MPa), the TRS and hardness of the MPEA binder-cemented carbides fabricated in this experiment were significantly improved, reaching 2741.5 MPa and 91.1 HRA, respectively. Compared with the pure cobalt binder-cemented carbides manufactured under the same conditions, the physical properties are better, and this binder has application prospects in the industrial field.