Novel Alternative Ni-Based Binder Systems for Hardmetals

Abstract

WC-Ni hardmetals, especially with the addition of Cr, are the first choice for wear parts in a corrosive environment. Despite Ni being studied as a metallic binder matrix in hardmetals for as long as Co, the mechanical properties achieved have consistently fallen behind those of their cobalt-containing counterparts. Due to the rapidly increasing demand for Co, its substitution is of increasing importance. In this study, various alloying elements that do not form strong carbides were systematically investigated as part of a binary Ni-based binder system for hardmetals. Solid and liquid phase sintering were compared by using field assisted sintering and a conventional SinterHIP furnace. The obtained hardmetals were analysed in terms of their microstructure, phases, sintering behaviour, and mechanical properties. The metals manganese, iron, and copper, as well as the metalloids silicon and germanium, were evaluated as additional binder constituents. Hardmetals with a binary Ni-based binder alloy were successfully prepared. The combination with Mn or Si showed the potential to significantly lower the necessary sintering temperature. In particular, Mn proved to be the most effective grain growth inhibitor among the investigated alloying elements.

1. Introduction

Hardmetals can be described as metal matrix compounds with the hard ceramic phase tungsten carbide (WC). The main task of the binder metal is to add toughness to the otherwise brittle ceramic. Over the course of hardmetal development, iron (Fe), cobalt (Co), and nickel (Ni) became the main contenders. Among these three metals, Co has the edge over the other two when comparing the achievable hardness–toughness ratio of the resulting hardmetals [1,2,3]. Various alloys of these three elements have been tested. A partial substitution of Co with Ni is possible without degrading the mechanical properties [4]. Combinations of all three elements, especially FeNiCo in a ratio of 40/40/20 wt.%, show good properties but have limited use [5,6]. Fe-based binder systems are still inferior in comparison, but new compositions are promising [7]. In the industry, WC-Co is by far the most important hardmetal. Advancing developments increased its performance further [8,9,10]. WC-Co is used in cutting, milling, mining, hard facings, wear parts, etc., in applications across the industry that require a hard yet tough material. The interest in the substitution of Co is increasing with the rising demand for other applications, especially for high-capacity Li-ion batteries. Besides WC-Co hardmetals, Ni as binder metal has some use because of its superior corrosion resistance [11,12,13]. Especially with additions of chromium (Cr) and molybdenum (Mo), it is used in environments that require high corrosion resistance [14,15,16], like processing of chemicals, pharmaceuticals, food, and many more. One of the main challenges when using Ni as binder metal is the promotion of WC grain growth during liquid phase sintering. Grain growth leads to a change in properties, e.g., a decrease in hardness. This can be suppressed by the addition of secondary carbides like VC, Cr₃C₂, and TaC [17,18,19,20,21]. Another effective way to inhibit excessive grain growth is the use of field assisted sintering techniques (FAST, also known as spark plasma sintering, SPS). This enables the sintering of hardmetals at comparably short times and at low temperatures in solid state, and thus diffusion and grain growth are greatly reduced [22,23,24]. In addition, the process is much faster than conventional vacuum sintering or SinterHIP. This is achieved by using mechanical pressure in combination with direct Joule heating of the conductive hardmetal powder via a current. While the necessary sintering temperature during liquid phase sintering of WC-Ni hardmetals are around 100 °C higher than for WC-Co (1350 °C), ca. 200 °C less is required to yield a fully densified body when FAST is used [25]. Next to the addition of carbide-forming elements like Cr or Mo to WC-Ni hardmetals, also investigations on non-carbide-forming elements that will be part of the metallic binder have been performed. Fe was studied in various shares in Ni-Fe binder systems. However, Ni is usually the smaller portion, making the binder essentially Fe-based. Although 50/50 (wt.%) mixtures of Fe and Ni as binder alloy provide good mechanical properties rivalling those of WC-Co hardmetals [26], smaller additions of Fe do not seem to be beneficial [27]. Limited work has been performed on Mn containing Ni-based binders. Here, an increase in high temperature flexural strength was shown [28]. The incorporation of small quantities of silicon can enhance the mechanical properties too [29,30]. Although the properties of Ni-Cu alloys are well-documented, there is an absence of studies on WC-Ni hardmetals containing Cu. However, investigations on conventional WC-Co hardmetals have revealed the potential use of Cu as a grain growth inhibitor when incorporated as part of the metallic binder [31,32,33]. So far, no systematic study has been performed. Therefore, in this study, a screening of potential alloying elements for binary Ni-based binder systems was carried out. The resulting compositions were prepared and sintered by conventional SinterHIP and FAST. These novel WC-Ni-based hardmetals were studied in terms of their microstructure, phase composition, and mechanical properties. To investigate the sintering behaviour of the hardmetals with the novel binder systems, dilatometric measurements and differential scanning calorimetry were carried out. The results are compared with thermodynamic calculations.

2. Materials and Methods

2.1. Materials and Composition

Six different compositions were developed and examined in this study: five WC-Ni-based alloys with the addition of Mn/Fe/Cu/Si/Ge and one reference with only Ni binder. To increase the effects of alloying, the binder amount was chosen to be at 20 vol.%, which is at the higher end of commercially available hardmetals. For pure WC-Ni, these 20 vol.% convert to a gravimetric share of 12.45 wt.%. Keeping the volumetric share constant across all compositions should ensure comparability. In contrast, the amount of a secondary binder constituent was defined by a constant atomic share. Here, the amount should be high enough to create distinct effects but low enough to avoid secondary phases like intermetallics, which are known to be brittle at low temperatures and can weaken mechanical properties. Thermodynamical calculations supported the selection process. They were carried out by using the CALPAHD method and were conducted with the software FactSage (v8.0, GGT-Technologies, Herzogenrath, Germany) in conjunction with the databases SGPS 2019 (v13.1) for pure substances and SGTE 2014 for intermetallic compounds and alloys (Scientific Group Thermodata Europe, Saint Sulpice, France). Based on these criteria, the share of the metallic elements Mn, Fe, and Cu was chosen to be 15 at.% and 7 at.% for the metalloids Si and Ge. The investigated compositions were designated as WC-NiX, with X denoting the respective alloying element: Mn/Fe/Cu/Si/Ge. The compositions of the corresponding hardmetal mixtures are presented in

Table 1. Composition of powders and theoretical densities (TD) of sintered WC-20 vol.% Ni_0.85X_0.15 samples.

Designation	Content/wt.%			TD/g/cm³
	WC	Ni	Mn/Fe/Cu/Si/Ge
WC-Ni	87.55	12.45	-	14.318
WC-NiMn	87.85	10.42	1.72	14.269
WC-NiFe	87.76	10.48	1.76	14.285
WC-NiCu	87.55	10.45	2.00	14.318
WC-NiSi	88.66	11.07	0.40	14.159
WC-NiGe	88.15	10.84	1.01	14.221

along with the theoretical densities calculated using the rule of mixture.

Only commercially available powders were used for this study. A WC powder with a medium particle size of 1.4 µm was chosen, as it represents a WC grain size present in many commercial hardmetal grades. The hardmetal should thus be comparable to commercial WC-Ni hardmetals and should offer mechanical properties in the range of a hardness of ca. 1300 HV30 and a fracture toughness of around 11 MPa*m^1/2. More details on the used starting powders are given in

Table 2. Details on the raw materials.

Powders	Average Particle Size/µm	Manufacturer
WC	1.4	H.C. Stark Tungsten GmbH, Goslar, Germany
Ni	2.5	Eurotungstene, Grenoble, France
Mn	10.0	GoodFellow, Cambridge, UK
Fe	1.6	BASF SE, Ludwigshafen am Rhein, Germany
Cu	0.9	CNPC Powder N.A. Inc., Vancouver, BC, Canada
Si	<10 µm	Elkem ASA, Oslo, Norway
Ge	<250 µm	HMW Hauner GmbH & Co. KG, Röttenbach, Germany

.

2.2. Powder Preparation and Sintering Process

For both studied sintering techniques, the same powder mixtures were used. To be suitable for the FAST process, the mixtures did not contain any organic binder. Hardmetal mixtures were prepared by dry mixing and subsequent wet milling in a planetary ball mill; 4 mm hardmetal balls with a powder-to-ball ratio of 1 to 10 were used. The powders were milled for 24 h at 70 rpm in n-heptane as a solvent and within a nitrogen atmosphere to prevent excessive oxidation. The n-heptane was removed using a vacuum furnace. A granulate was produced by sieving (315 µm mesh) the dried powder–ball mixture. One part of the powder mixtures was consolidated via FAST (H-HP D 25, FCT Systeme GmbH, Frankenblick, Germany). This technology uses a direct current to heat the conductive powder by the Joule losses as well as the tools made from graphite. To protect the die from excessive tool wear and to provide a good seal, the cavity was lined with graphite foil in between die and punch as well as punch and powder. The powder was manually filled into the cylindrical dies (inner Ø: 20 mm). Samples of 5 mm hight were produced. This ensured having enough distance from edge-induced inhomogeneities caused by temperature and pressure gradients as well as absorbed carbon from the graphite foil. In addition, a graphite felt around the tool for insulation was used to further minimize temperature gradients. The powder was pre-compacted within the tool at 10 kN. The temperature was in-line controlled via a pyrometer. The measurement area was located at the bottom of the blind hole inside the top punch. To minimize the process time without increasing temperature gradients too much, a maximum heating rate of 300 K/min was used. Up to 600 °C, a lower heating rate of 100 K/min was used to ensure degassing of adsorbed gases and water. At 50 K below the maximum sintering temperature the heating rate was lowered to 100 K/min within 30 s to prevent overshoot. Simultaneously, the piston force was increased from the initial 10 kN to a maximum of 25 kN (32 MPa to 80 MPa for Ø 20 mm die). The dwell was set at 5 min for all trials. Two maximum temperatures of 1100 °C and 1140 °C were studied. Sintering was conducted in vacuum. Among other parameters, temperature, displacement, and chamber pressure were recorded. Due to the short sintering time and temperatures as well as the absence of a liquid phase, it was expected that a thermodynamically stable state could not be achieved. For comparison, the CIPed green parts were heat treated in a SinterHIP furnace. In order to form green parts for the sintering process, the powders were cold isostatically pressed at 300 MPa to a near-cylindrical shape. Due to the lack of an organic binder, no debindering stage was necessary. For WC-Ni hardmetals, higher sintering temperatures are necessary than for WC-Co to achieve liquid phase sintering. A temperature of 1450 °C was used. The samples were sintered in one run and the same graphite crucible. Heating and cooling were performed at a rate of 10 K/min. During dwell at maximum temperature, the pressure was adjusted to 0.01 MPa with a flow of Ar to prevent excessive Mn evaporation due to its relatively high vapour pressure at higher temperatures. Subsequently, Ar pressure was increased to 10 MPa to reduce residual porosity during the “HIP” stage of the SinterHIP process.

2.3. Sample Preparation and Characterization Methods

The FAST sintered samples were first sandblasted and ultrasonically cleaned to remove dust from the surface. Then, together with the SinterHIP treated samples, the density was measured according to ISO 3369 (Archimedes principle). After cutting them in half, the samples were embedded into resin and polished to a mirror finish using diamond-based suspensions. This included microstructural examinations by field emission scanning electron microscopy (FE-SEM; ULTRA 55, Carl Zeiss Microscopy GmbH, Jena, Germany) and energy dispersive X-ray spectroscopy (EDS; X-MaxN 80, Oxford Instruments plc, Abingdon, UK) with the detector being attached to the FE-SEM. Phase analysis was carried out by X-ray diffraction (XRD; D8 ADVANCE, Bruker Corp., Billerica, MA, USA) using Cu Kα radiation. The linear intercept method was used according to ISO 4499-02 to gain information about the impact of alloying the Ni binder on WC grain size. Measured values correlate with grain size. The arithmetic mean as well as the 10%, 50% (median), and 90% quantile were determined. Carbon content was measured via the analysis of gases from the combustion of the powdered material in pure oxygen (WC600, Leco, St. Joseph, MI, USA). As a reference material, high-purity WC (ECRM No. 783-1, BAS Ltd., Middlesbrough, UK) with a defined C content of 6.188 ± 0.013 wt.% was used. For thermophysical examinations, the CIPed powder was used. Dilatometric measurements were carried out to study the shrinkage during the liquid phase sintering process using a push rod dilatometer (DIL 402 C/7/G, NETZSCH Gerätebau GmbH, Selb, Germany) in a temperature range between RT and 1500 °C according to DIN 51045-3. Complimentary differential scanning calorimetry (DSC, DSC404F1, NETZSCH Gerätebau GmbH, Selb, Germany) was conducted to examine the formation of the liquid phase. An uncalibrated measurement was used. Therefore, heat flows cannot be compared directly. The hardness of the hardmetals was measured via the Vickers indentation method (HV30; ISO 3878). By measuring the length of the cracks that propagate from the corners of the indentation, the fracture toughness was calculated according to ISO 28079.

3. Results and Discussion

3.1. Thermodynamic Calculations

For choosing suitable alloying elements and to evaluate their content within the Ni-binder, the corresponding phase diagrams were calculated by means of CALPHAD. For further consideration, the amount of the studied alloying elements was fixed to 15 at.% for Mn, Cu, and Fe and 7 at.% for Si and Ge. For the metalloids, the share was chosen to be lower, to avoid the formation of intermetallic phases. The phase diagrams with 20 vol.% binder content are shown in Figure 1. All calculations were performed for the thermodynamically stable state. This includes the possible formation of the metastable cementite (Fe₃C) in WC-NiFe. Due to the low carbon content of the melt above and the binder within the two-phase area as well as the high solubility of carbon in the fcc lattice, the formation of either primary or secondary cementite is unlikely. Without a phase transformation to bcc accompanied with a drop in solubility for carbon, also no ternary cementite or martensite is expected to form. When comparing the calculated phase diagrams, it is apparent that the addition of secondary elements to the Ni binder influences the two-phase area (just WC and Ni-based binder) in multiple ways. Different solubilities for W and C will shift the two-phase area to different C contents. Its relative position towards the theoretical carbon content that is calculated for a 1:1 stochiometric WC with the respective binder is crucial. Ideally, it should be within the two-phase area, which is the case only for the compositions with Mn. However, in practice, the carbon content is also greatly influenced by the starting powders (carbon and oxide content) and sintering processes (carbon loss due to reduction of oxides and carbon pick-up). Especially for the Mn and Ge containing compositions, the solidus temperatures decrease significantly with increasing carbon content within the range of the two-phase area. Adjusting the carbon content to the upper end of the two-phase area can therefore help with densification without the need for an increased sintering temperature. In contrast, the composition with Cu addition only shows minor differences in the solidus temperatures with varying C content. The width of the two-phase area is also influenced by the alloying elements. Mn, Fe, and Si narrow it, whereas Cu and Ge do not influence the width greatly in the considered temperature range. A wider window is beneficial since less effort is required to adjust the carbon content during the hardmetal production. The needed sintering temperature to consolidate WC-Ni hardmetal is usually about 100 °C higher than for WC-Co. Higher sintering temperatures promote grain growth. Theoretical calculations show that especially the addition of Mn and Ge shift the solidus line to significantly lower temperatures, which can potentially reduce the total time and energy consumption of the sintering process. Furthermore, no additional phases like intermetallics are predicted to form with the studied additions. Figure 1 also shows the measured C content of the liquid phase sintered samples (green line). For the compositions with Fe, Cu, and Ge, as well as for the pure Ni binder, the C content is within the two-phase area. For WC-NiSi only above ca. 1100 °C, a two-phase microstructure is stable at the measured C content. Upon cooling, a mixed carbide (eta carbide, M₆C) can form. However, in solid state the diffusion is limited below this temperature for the bigger atoms W and Ni and the equilibrium state might not be reached during the FAST process. WC-NiMn shows the highest carbon deficit indicating the formation of eta carbide.

3.2. Thermophysical Properties

To study the influence of the different additions to WC-Ni hardmetals on the sintering behaviour, dilatometric tests were carried out. Since powders were pressed via CIP and therefore did not have a geometrically defined shape, exact measurements of the green density were not possible. This makes the maximum value of the relative shrinkage not comparable. More importantly are the characteristics in the course of the sintering process. The relative shrinkage curves show that all compositions already start to densify at temperatures above 800 °C (Figure 2a). Up to around 1000 °C, they all share a common solid state sintering behaviour. Most notably are the differences in the maximum relative shrinkage rate (Figure 2b). These can be explained with the different temperatures at which a liquid phase occurs. The two areas of solid and liquid phase sintering can be differentiated by two characteristic extrema. The separation of the two processes is especially pronounced for the WC-Ni reference. The maximum shrinkage rate is achieved at 1242 °C during the solid state sintering phase and has slowed down significantly before reaching the temperature of liquid phase formation above 1400 °C. In the case of WC-Ni this is even accompanied by a sudden expansion due to the lower density of the liquid. On the other hand, liquid phase forms for WC-NiMn during the most active phase of solid state sintering, which leads to the formation of only one significant peak and the highest densification rate. The other compositions show states in which the extrema overlap only partially. During solid state sintering, the kinetic processes that lead to densification of the powder are significantly slower in comparison to liquid phase sintering. This can be directly observed in the relative shrinkage rate graphs from WC-NiFe. Despite most of the shrinkage already appeared during solid state sintering, the maximum rate is achieved upon formation of liquid phase. However, in solid state, grain growth is not as pronounced. In the liquid phase, WC will dissolve and precipitate at a higher rate. Even abnormal grain growth is possible. It is well known that sintering behaviour of hardmetals strongly depends on the C content [34]. This becomes clear when looking at the phase diagram of WC-Ni. More than 80 K separate the lowest solidus temperature from the highest above the two-phase area of WC and Ni. The addition of elements to the hardmetal composition will shift the relative position of the theoretical C content and the two-phase area. For WC-Ni, it is considerably higher and outside the two-phase area. With Mn, the line is shifted into it. Taking into consideration a similar yet unknown carbon loss for all compositions (due to the reduction of surface oxides), the formation of a liquid phase should be shifted towards higher temperatures for all compositions. The composition with Mn densifies at the lowest temperature because even the highest solidus temperature is among the lowest across all compositions. Only WC-NiGe forms a liquid at a lower temperature at the higher end of the C content. Looking at the plots of the relative shrinkage rate, it is worth noting that the maximum shrinkage rate of the solid phase sintering stage of the Si containing compositions is at a lower temperature than for pure WC-Ni. However, Cu and Fe seem to increase this temperature.

The temperature at which a liquid phase forms can be estimated by using the onset method on the dilatometric diagrams. However, when the effects of solid and liquid state sintering strongly overlap, as with WC-NiMn, this estimation becomes difficult. The relevant processes must be known for the correct interpretation of dilatometric results and determination of onset temperatures. Therefore, DSC measurements were carried out (Figure 3a). Here, the findings are in line with theoretical calculations and shrinkage behaviour. Pure WC-Ni shows the highest onset at 1422 °C and WC-NiMn the lowest at 1275 °C, while the other compositions are in between. At which temperatures the endothermic phase transition starts and ends as well as the maximum value of the received signal depends on the transition area between the solidus and liquidus line. For WC-Ni, the endothermic peak is sharp, whereas for WC-NiMn, the peak is flat and wide and shifted towards lower temperatures.

Figure 3b shows a comparison of temperatures at which a liquid phase forms. Temperatures were derived from theoretical calculations, dilatometric and DSC measurements. Theoretical values were taken from the CALPHAD calculated phase diagrams at the measured C content. For the dilatometry measurements, the onset temperature had to be estimated from the relative shrinkage rate graphs (Figure 2b). Here, it is crucial to take into account the effects that take place during sintering, as described before for choosing the correct onset. The areas that are associated with liquid phase formation are indicated by dots. For WC-NiMn, no distinction between the solid and liquid phase sinter processes from the dilatometric data was possible. The kink at about 1345 °C and the following increased incline in the graph of the relative shrinkage rate might be an indication for the end of liquid phase formation. Considering the data from theoretical calculations and DSC leads to the conclusion that for WC-NiMn during active solid state sintering process, a liquid phase forms, and the associated effects increase the overall shrinkage rate. Therefore, the onset temperature was derived from the relative shrinkage plot. Figure 3b illustrates that especially Mn but also Ge and Si can significantly reduce the temperatures of liquid phase formation. For Fe and Cu addition, the solidus temperatures should be slightly lower as well. However, a distinct advantage for sintering at significantly lower temperatures is not expected.

3.3. FAST Densification Behaviour

One of the advantages of FAST is the ability to directly monitor the densification behaviour. Also, the pressure inside the vessel, which is an indicator for degassing and reduction processes, can be analysed easily. In Figure 4, both the displacement (shrinkage) and the final relative density for the two used sintering temperatures are shown. For a better comparison, the displacement has been set to zero at 600 °C. At this temperature, degassing of adsorbed gases and water is finished, and no sintering processes have started yet. At a first glance, the differences in maximum displacements of the different compositions at 1140 °C becomes obvious (Figure 4a). This could be attributed to a different magnitude of residual porosity or the formation of secondary phases. However, it is also dependent on the green density of the pre-compacted powder, which is different for each composition and depends on the manual powder-feeding process. More important are the time and temperature at which no further densification takes place. Especially, the composition with Mn stands out, but also the compositions with Si and Ge, which reach the plateau earlier as compared to the pure Ni binder. The addition of Fe, on the other hand, leads to a delayed densification.

The influence of the alloying elements on relative densities (RD) is evident and supports the findings from the displacement graphs when comparing them at a sintering temperature of 1100 °C (Figure 4b). The addition of Si and Ge and to a certain extent Mn leads to a higher RD than pure Ni. This is in line with the calculated lower solidus temperatures. In case of Mn, already at 1100 °C no further compaction takes place. This could indicate the formation of a ternary phase with a lower density and/or the development of gas due to reduction processes, which could lead to porosity. Only the WC-NiFe sample sintered at 1100 °C has a lower RD compared to WC-Ni. Even at a higher sintering temperature of 1140 °C, the displacement for WC-NiFe takes the most time to reach a plateau. This densification behaviour correlates with the results from dilatometric measurements (Figure 2b). During the solid state phase the maximum relative shrinkage rate is lower and shifted towards a higher temperature in comparison to WC-Ni. The RD above 100% of WC-NiGe can be explained by a lower share of Ge in the sintered sample compared to the target value. Subsequently performed XRD measurements of the Ge powder revealed a significant amount of GeO₂. During sintering, the oxide is reduced by carbon. Since Ge is part of the binder, the lower amount of the alloyed metallic binder leads to a higher share of WC, increasing the overall density above 100%.

3.4. Phase Analysis and Microstructure

Since elemental powders for the metallic binder were used, it is necessary to verify successful mixing. This is especially important for the samples produced by FAST, since no liquid phase is formed here. According to the calculated phase diagrams, a two-phase area of WC and the metallic binder exists for all compositions. This is confirmed by XRD measurements (Figure 5). Only in WC-NiMn a third phase can be detected after SinterHIP treatment. It identifies as the mixed carbide M₆C, also known as eta carbide, which forms at lower C contents outside the two-phase area and can reduce the toughness of the hardmetal. The detected pattern correlates quite well with the peaks of Ni₂W₄C. Also, no cementite could be detected in WC-NiFe. Furthermore, it is notable that the peaks of the Ni-based binder phases are slightly shifted depending on the alloying element. This is especially pronounced in the SinterHIP treated samples. In diffractograms, positions of peaks for pure Ni (fcc) are given for reference. Even the composition without the addition of a secondary binder constituent shows a shift in relation to pure Ni of the binder peaks due to W and C being dissolved.

EDS mappings of the FAST processed samples, shown in Figure 6, support the XRD results. Elements are evenly distributed across the binder areas. No clustering of the alloying elements could be detected. A mapping with sufficient resolution and contrast could not be performed for WC-NiSi. Emission peaks of Kα₁ of Si and Mα₁ of W are separated only by 34 eV, which is below the resolution of the used detector. Even when trying to identify W by other electron transitions, the weak signal from Si would be masked by the intense signal of W due to the comparably low concentration of Si. Exciting higher energetic transitions with higher acceleration voltages leads to a decrease in spatial resolution due to an increase in the exited volume, which is crucial in detecting small inhomogeneities in the binder. Therefore, W and Si could not be distinguished within the binder area and hard phase using EDS.

The optical micrographs show a fully dense body with a porosity better than A02 according to ISO 4499-4 for all compositions. SEM micrographs, as shown in Figure 7, reveal a quite uniform microstructure across all FAST sintered samples. The WC grains are irregularly (unfaceted) shaped, and their size is in the range of the powder particles with a high number of grains smaller than 1 µm. Mn and Si containing compositions show third phases, which appear darker in the BSE images than the other phases. This indicates the presence of a higher share of elements with a lower atomic number. These third phases cannot be found in the SinterHIP treated samples.

Due to the presence of a liquid phase during SinterHIP treatment and therefore facilitated diffusion processes that promote the dissolution and precipitation of WC within the liquid binder, WC grains in all samples are more faceted and show the angularly shaped morphology typical of liquid phase sintered hardmetals. This process also leads to the dissolution of the finest grains and grain growth during liquid phase sintering. These effects are especially pronounced in the non-alloyed WC-Ni. Grains are bigger for the SinterHIP treated samples in comparison to those prepared by FAST. Alloying leads to a grain growth inhibition in comparison to pure WC-Ni for the liquid phase sintered samples.

Table 3. Results of linear intercept measurements of the SinterHIP treated samples.

Sample	dmean, µm	d10, µm	d50, µm	d90, µm
WC-Ni	0.80 ± 0.54	0.25	0.68	1.52
WC-NiMn	0.46 ± 0.29	0.17	0.40	0.81
WC-NiFe	0.59 ± 0.38	0.19	0.50	1.08
WC-NiCu	0.61 ± 0.44	0.19	0.52	1.13
WC-NiSi	0.52 ± 0.33	0.19	0.44	0.98
WC-NiGe	0.64 ± 0.41	0.21	0.56	1.15

summarizes the linear intercept measurements of WC grains, which correlate with grain size. All samples show a monomodal grain size distribution. The WC-NiMn samples show the smallest grain size. This could not only be related to grain growth inhibitive properties of Mn itself but also the low carbon content in relation to the two-phase area, which even leads to the formation of eta carbide. WC-Ni has the biggest grains. The other compositions are in between these extremes, with Cu and Ge acting as the least effective grain growth inhibitors.

Ternary phases found in WC-NiSi and -NiMn were analysed by EDS (Figure 8 and Figure 9). The measurements revealed a high oxygen concentration in these third phases. They are most likely oxidic species of the respective element, with MnO and SiO₂ being the most stable oxides in the presence of C and CO at the sintering temperatures used. Despite Si being not distinguishable in the binder area or the hard phase from W in the low amount used, it can be clearly detected in the oxides due to the absence of W and the higher concentration of Si.

3.5. Mechanical Properties

Hardness and toughness were measured via the Vickers indenter method and crack length measurement. As seen in Figure 10, alloying WC-Ni leads to a decrease in tough-ness and an increase in hardness regardless of the sintering method. This could be at-tributed to solid solution strengthening of the binder, which lowers its toughness. For samples treated by SinterHIP, the grain growth inhibiting effect could be a factor too. A finer grained hardmetal is usually harder but has a lower toughness if all other factors are identical [9]. This also explains the different hardness-to-toughness ratio between the two sintering technologies. The liquid phase sintered samples have coarser grains. Only WC-NiMn is an exception. The FAST sintered WC-NiMn sample is less hard but slightly tougher in comparison to the SinterHIP sample. Lower hardness can be attributed to the formation of Mn oxides. The WC-NiSi sample treated by SinterHIP shows a significant increase in hardness in comparison to the other compositions (except NiMn). This can be attributed to the second smallest WC grain size and is in line with the typical hardness–toughness trade-off associated with finer WC grains.

Despite not having achieved the target value regarding the Ge content, the liquid phase sintered WC-NiGe also shows an increase in hardness. However, when put into relation to the fracture toughness the ratio is worse compared to the non-alloyed WC-Ni. The error of the measured toughness increases with rising values because the crack lengths are decreasing, and slight variations lead to significantly different readings. The highest fracture toughness of 20.6 MPa*m^1/2 was thus found for the pure WC-Ni produced by SinterHIP, and the highest hardness of 1377 HV30 was found for the WC-NiFe sample produced by FAST.

4. Conclusions

In this work, hardmetals with alternative Ni-based binder systems were developed, and their microstructure as well as properties were investigated. They were consolidated via conventional liquid phase sintering in a SinterHIP furnace at 1450 °C as well as by field assisted sintering, which made it possible to densify the powders with direct Joule heating and mechanical pressure in solid state below 1150 °C. The results were compared regarding microstructural features, phase composition, and mechanical properties. The influences on sintering behaviour and properties by adding Mn/Fe/Cu/Si/Ge to WC-Ni hardmetal were examined for both processes supported by dilatometric and calorimetric tests. The findings were compared to theoretical calculations obtained by the CALPHAD method. The following conclusions can be drawn:

• WC-Ni + Mn/Fe/Cu/Si/Ge can be consolidated by SinterHIP as well as FAST with the formation of a two-phase microstructure with a single-phase alloyed binder and the hard phase WC.
• Mn and Si oxides were found in the FAST prepared samples, which could not be reduced due to the short sintering time and low sintering temperature.
• Theoretical calculations predicted a significant decrease in solidus temperatures in the C range of the two-phase area in comparison to the pure Ni binder, especially with the addition of Mn, Si, and Ge. This could be confirmed by dilatometric and DSC measurements.
• Alloying WC-Ni leads to a shift of the two-phase area relative to the C content of the compositions. This was most pronounced in WC-NiMn and leads to the formation of the mixed carbide eta carbide in the SinterHIP sample.
• While the WC grain size for the FAST consolidated compositions was quite uniform, those treated by SinterHIP with the formation of a liquid phase revealed a certain grain growth inhibiting effect with the addition of the alloying elements to WC-Ni. This was most obvious in the Mn and Si containing samples.
• The overall bigger WC grain size of the SinterHIP treated samples leads to a higher toughness and lower hardness in comparison to samples prepared by FAST, with the exception of WC-NiMn due to the formation of eta carbide.
• With the addition of Mn and Si, hardness increased significantly for the SinterHIP treated samples.

Despite the formation of oxidic inclusions during FAST consolidation, the addition of Mn and Si shows the potential of improving the hardness of liquid phase sintered WC-Ni hardmetals. Additionally, sintering temperatures can be significantly decreased with the addition of these two elements. This reduces grain growth and last but not least overall energy consumption. Future investigations will focus on the influence and the adjustment of the C content, variations in the shares of alloying elements, their combination, as well as lower binder contents.