Process Studies on the W-C-Ti System Using a High-Throughput Laser-Based Additive Manufacturing Approach

Abstract

Conventional WC-Co hard metals have proven to be difficult to manufacture by means of laser powder bed fusion (PBF-LB), resulting in residual pores, crack formation, foreign phase formation, and the inhomogeneous growth of the carbide phase. Alternative compositions such as the W-C-Ti system presented in this study need to be investigated. Through the employment of a high-throughput screening approach, 11 alloy compositions were investigated to determine the influence of the carbon content and tungsten–titanium ratios on microstructure formation and basic mechanical properties. Two screenings were conducted, with one varying the carbon content (10–35 at.%) and the other adjusting the W/Ti ratios (10:90 to 60:40 at.%). Microstructural analyses using scanning electron microscopy (SEM), X-ray diffraction (XRD), and hardness measurements provided insights into phase formation, grain distribution, and mechanical properties. The results showed that increasing the carbon content significantly enhanced the hardness (from 681 HV (10 at.% C) to 1898 HV (35 at.% C)) due to higher δ-(Ti,W)C_1−x carbide phase fractions. Alloys with a higher tungsten content exhibited finer microstructures and an improved crack resistance while maintaining a high hardness (1900–2100 HV). This study identified an alloy with 32.5 at.% W, 32.5 at.% Ti, and 35 at.% C as a promising candidate for further investigation, with properties similar to those of a conventional WC-Co hard metal.

1. Introduction

Due to their excellent properties in terms of Young’s modulus, hardness, wear resistance, and sufficient fracture toughness, WC-Co hard metals are suitable for a wide range of applications, such as cutting and punching tools, deep drawing and wire drawing applications, wear parts such as milling balls, and machine parts such as rollers and tool carriers.

Cobalt has been established as the primary binder material due to its relative low melting point enabling liquid-phase sintering, its extraordinary wetting behaviour towards tungsten carbide, and its ductility, which leads to the sufficient fracture toughness of a specific component. These excellent properties make the conventional production of hard metal components via powder metallurgical processes with subsequent mechanical finishing a cost- and labour-intensive process. However, a study from Dedalus Consulting from 2022 [1] showed that 53% of cutting materials are covered by cemented carbides and 27% by super hard cutting materials such as ceramics, PcBN, and PCD, which have a similar manufacturing process. More than half of the cutting materials used in the world are therefore composed of tungsten and cobalt, which can both be found on the EU list of critical raw materials (CRMs) [2]. In addition, both pose health risks, with special regard to cobalt [3]. To address the criticality of the most common cutting materials and also the increasing necessity for specialized cutting tools, the adoption of additive manufacturing (AM) shows high potential.

Due to layer-by-layer manufacturing in AM, where only the required amount of material is used to manufacture the desired component, postprocessing and thus waste material can be significantly reduced.

However, WC-Co has been shown to be difficult to produce by means of direct-melting AM techniques. In exemplary publications on the laser-based AM of WC-Co hard metals, Uhlmann et al. [4], Fortunato et al. [5], Genilke et al. [6], and Kumar et al. [7] all produced samples with high porosity, cracks, undesirable phases, and inhomogeneous carbide grain size distributions and thus improper material properties. In particular, the undesired η-phase and the inhomogeneous growth of the WC particles led to a weak fracture strength and toughness. Only Khmyrov et al. [8] were able to produce crack-free samples with minimal brittle phases while using very high binder contents of up to 75% Co. All of the above-mentioned authors, as well as several more not mentioned specifically here, came to the conclusion that the very high laser energy required to melt the composite powders and achieve a suitable densification ultimately leads to the substantial dissolution of the carbide phase and thus the formation of brittle phases and the inhomogeneous formation of the hard phase. In addition to the geometrical freedom of AM processes and the possible reduction in material waste, the very high temperatures that can be achieved during laser melting make the PBF-LB process very interesting for producing and investigating previously almost unthinkable combinations of high-melting-point metals and hard-phase-forming elements, thus enabling new material combinations to be found with reduced CRMs but with sufficient properties.

In this study, the material system W-C-Ti was investigated through the application of a high-throughput screening method that first provides insight into the microstructure formation and basic mechanical properties of different material combinations while using small amounts of powder without needing PBF-LB-specific powder properties. The aim was to find a novel material system for PBF-LB WC-based composites without Co and to also achieve a possible reduction in W. Ti was chosen in this study because it is not considered a CRM, it forms a carbide with a desirable hardness [9], and it stabilizes the mono-carbide by suppressing W₂C formation [10,11]. A deeper look into the W-C-Ti system with regard to constitution and phase formation was carried out in 1984 by Uhrenius [10] and in 1999 by Haldar et al. [11]; however, the authors did not specifically look into material properties or the application of the resulting materials. Hofmeister et al. [12] investigated the fabrication of cast carbide–metal composites in the W-Ti-C system via laser powder deposition. They produced near-net-shape components consisting of a eutectic microstructure of finely distributed (Ti,W)C dendrites in a (W,Ti) metal matrix. Their work confirmed the formation of hypereutectic microstructures with Vickers hardness values ranging from 1100 to 1300 HV, emphasizing the potential of laser processing for refractory systems where conventional sintering may be impractical. Another study published by Reva et al. [13] found that ultrafine WC and TiC powders synthesized mechanochemically produced Ti-W-C hard alloys for cutting material with enhanced mechanical properties, such as plasticity, toughness, and an extended cutting tool life. Further comprehensive studies specifically targeting the W-C-Ti system are scarce, which underlines the need for investigations such as the one presented in this paper to explore the potential of these materials.

With the two carried out screenings, 11 different compositions were investigated, analyzing the influence of the carbon content and the tungsten–titanium ratio. The characterization of the produced samples was carried out by means of microstructural examinations using scanning electron microscopy (SEM), a quantitative image analysis, Vickers hardness measurements, and X-ray diffraction (XRD) phase analyses.

2. Materials and Methods

2.1. Alloy Selection to Investigate Carbon Content and Tungsten–Titanium Ratio Influence

The influence of the carbon content and the tungsten–titanium ratio on microstructure formation was investigated using two screenings with six samples each. Figure 1 and Figure 2 display the ternary W-C-Ti system in the liquidus projection (Figure 1) and the isothermal section at 1500 °C (Figure 2) according to [11].

The liquidus projection shows that five areas exist, from which five different possible primary solidifications emerge, namely C_Graphite, ε-WC, δ-(Ti,W)C_1−x, γ-W₂C, and β-(Ti,W). The two promising primary crystallization areas are highlighted in Figure 2.

The phase equilibria are displayed in the isothermal section. The aim of the present study was to find a favourable two-phase composition providing a carbide and a metallic phase. To achieve this, all the investigated alloys were located in the two-phase area β-(Ti,W) and δ-(Ti,W)C_1−x (Figure 2).

In the first screening (vertical dots in Figure 1 and Figure 2), the influence of the carbon content on the hardness and phase formation was investigated. A fixed W/Ti ratio of 50:50 was used, and the carbon content was varied from 10 at.% to 35 at.% (

Table 1. Investigated compositions used to analyze the influence of carbon content.

Sample	W at.%|wt.%	C at.%|wt.%	Ti at.%|wt.%
10 at.% C	45.0|78.4	10.0|1.1	45.0|20.4
15 at.% C	42.5|77.9	15.0|1.8	42.5|20.3
20 at.% C	40.0|77.3	20.0|2.5	40.0|20.1
25 at.% C	37.5|76.7	25.0|3.3	37.5|20.0
30 at.% C	35.0|76.0	30.0|4.3	35.0|19.8
35 at.% C	32.5|75.2	35.0|5.3	32.5|19.6

). The targeted hardness range was 1000–2000 HV, and the phase composition included one carbide and one metallic phase. The double-carbide γ-W₂C was an undesired phase in this investigation due to its brittleness, typical of conventional WC-Co hard metals.

The second screening was conducted to investigate the influence of the W/Ti ratio. Six alloys with a constant carbon content and a varying W/Ti ratio were selected for this purpose. An amount of 35 at.% of carbon was chosen after the evaluation of Screening 1 based on the achieved hardness and microstructure. The W/Ti ratio was changed from 10:90 to 60:40 in steps of 10 to compositions with higher tungsten contents (

Table 2. Investigated compositions used to analyze the influence of the W/Ti ratio for fixed carbon contents.

Sample	W at.%|wt.%	C at.%|wt.%	Ti at.%|wt.%
W/Ti 10:90	6.5|27.06	35.0|9.52	58.5|63.42
W/Ti 20:80	13.0|45.10	35.0|7.93	52.0|46.97
W/Ti 30:70	19.5|57.98	35.0|6.80	45.5|35.22
W/Ti 40:60	26.0|67.63	35.0|5.29	39.0|26.42
W/Ti 50:50	32.5|75.20	35.0|5.30	32.5|19.60
W/Ti 60:40	39.0|81.15	35.0|4.76	26.0|14.09

). Since the undesirable phase γ-W₂C is formed at higher tungsten contents (Figure 2), the tungsten content was not increased further than 39 at.%.

2.2. High-Throughput Screening

For a time- and material-efficient investigation of different new material compositions in laser-based additive manufacturing, a high-throughput screening workflow was established. This approach was utilized for the present study. The process consists of five steps that include selecting the compositions, blending the starting powders, compacting the powder mixture into a handleable format, performing laser exposure with different parameters, and subsequently preparing and analyzing. The process is schematically shown in Figure 3. The alloy selection step is a prerequisite for the following steps and is described in the previous section.

The main focus of the high-throughput screening approach is to develop a first idea of the melting behaviour and the microstructure formation of experimental compositions while only having minimal powder quantities available that do not always have the necessary characteristics (sphericity, flowability) for a conventional PBF-LB process. Previous studies have shown comparable microstructure formation in the conventional layer-wise PBF-LB and the single-layer high-throughput screening approach utilized in this study.

2.2.1. Powder Blending

A homogeneous powder blend is an important prerequisite for the high-throughput approach to ensure a precise composition for all laser-molten samples. To ensure this, the initial powders (see

Table 3. Powders used for the experiments, manufacturer, grainsize and purity.

Type	Name	Manufacturer	Grain Size	Purity
W	W0.7 µm	Wolfram Bergbau- und Hütten AG (Mittersill, Austria)	0.7–0.8 µm (FSSS)	99.0%
WC	Tungsten Carbide DS100	Global Tungsten & Powders (Towanda, PA, USA)	1.0–1.2 µm (FSSS)	99.0%
Ti	Amperit 155.086	Höganäs Germany GmbH (Goslar, Germany)	≤63 µm	98.8%
TiC	STD120	H.C. Starck Surface Technology and Ceramic Powders GmbH (Munich, Germany)	1.0–1.5 µm (FSSS)	98.5%

) were first weighed according to the desired composition and then mixed with ethanol in a BEVS 2501/1 laboratory mixer (BEVS Industrial Co., Ltd., Guangzhou City, China) at 2000 rpm for 15 min with a dispersing disc. The homogeneous slurry was further dried overnight in a Nabertherm TR-062-N furnace (Nabertherm GmbH, Lilienthal, Germany) at 80 °C in an ambient atmosphere. SEM images of the irregularly shaped initial powders are presented in Figure 4a–d. For the high-throughput approach, spherical powders with a good flowability, as needed for the conventional PBF-LB process, are not necessary due to the mechanical compaction in the next step.

2.2.2. Compacting

After drying, the homogeneous powder blend was transferred into a 30 mm floating die pressing tool and subsequently compacted down to a sample height of 5.5 mm. The compaction was carried out in a Fontijne TP400 table press (Fontijne Presses b.v., Rotterdam, The Netherlands) at a 100 kN force (140 MPa) for 30 s in an ambient atmosphere. After pressing, the powder compacts were removed from the die and dried again overnight at 80 °C in a Nabertherm TR-062-N furnace (Nabertherm GmbH, Lilienthal, Germany) in an ambient atmosphere to ensure no residual moisture was present in the powder during laser melting. The compaction step was necessary in this approach to allow the powder to be handled and transferred to the process chamber. In pre-trials to this study, the minimum compacting force was determined to obtain handleable powder tablets without diverting too much from the loose powder bed characteristics in the PBF-LB process.

2.2.3. Laser Melting

The compacted and dried powder tablets were then placed in an inert gas chamber equipped with a laser-transparent window under an argon atmosphere at ~50 ppm residual O₂. The laser-melting process was afterwards carried out with a Trumpf TruFiber 1000 fiber Laser (TRUMPF SE & Co. KG, Ditzingen, Germany) and a Scanlab IntelliScan 30 scanner system (SCANLAB GmbH, Puchheim, Germany). For the present study, individual fields of 3.5 × 3.5 mm were exposed to the laser beam on the powder surface at linear energy densities varying from 0.38 to 2.5 J/mm. The 3.5 × 3.5 mm fields were further filled with laser lines with a hatch distance resulting in 20% track overlap. An exemplary representation of a remelted powder compact with an overlay of the applied laser parameters for each sample is shown in Figure 5.

In the present study, only the samples with the highest linear energy density (2.5 J/mm) were investigated. Those samples were remelted with a laser power (P_L) of 250 W and a scanning speed (v_s) of 100 mm/s. The samples were chosen because after the first inspection, they provided the largest molten volume and thus possessed the most suitable dimensions for characterization.

2.3. Characterization

For characterization, the 3.5 × 3.5 mm sample pieces were carefully separated from the powder compact and embedded in epoxy resin (Struers EpoFix, Struers GmbH, Willich, Gemany). The cross-section orientation was always perpendicular to the laser scanning direction. The sample preparation was carried out according to materialographic standards for hard and brittle materials with polishing cloths and diamond slurries with grain sizes from 9 µm down to 0.25 µm. A contrasting process (e.g., materialographic etching) was not necessary for the samples investigated in this study. Microstructure characterization was carried out using a ZEISS Sigma 300 VP field-emission scanning electron microscope (FE-SEM, Carl Zeiss Microscopy GmbH, Oberkochen, Germany) and backscattered electron contrast (BSE) to visualize the different phases in material contrast. A quantitative image analysis was performed on the SEM images using ZEISS Zen core (version 3.8, Carl Zeiss Microscopy GmbH, Oberkochen, Germany) software with machine learning-assisted phase segmentation (ZEISS ZEN Intellesis version 3.8, Carl Zeiss Microscopy GmbH, Oberkochen, Germany) to quantify the phase fractions. Each pixel in the image was assigned to either the carbide or the metallic phase, and the phase fraction was calculated based on the percentage of the total pixels of each class. The hardness was measured with Vickers in accordance with the current DIN EN ISO 6507-1 [13] on an EMCO-Test Durascan 70 micro-hardness tester (EMCO-TEST Prüfmaschinen GmbH, Kuchl, Austria). Due to the sample size, the Vickers method used was HV 0.5, corresponding to a load of 4.903 N. Phase identification was achieved with a Seifert Sun XRD 3003 fast in situ X-ray diffraction machine (XRD, Waygate Technologies, Hürth, Germany) using a Cu tube with 35 kV and 50 mA with a 15–135° 2-theta range.

3. Results

3.1. Screening 1—Influence of Carbon Content

Figure 6a–f shows the representative microstructures of the first six compositions with increasing carbon content (Table 1). Based on the primary crystallization, the six compositions can be divided into two groups, where the samples with 10, 15, or 20 at.% C showed a bright phase as the primary phase and the compositions with 25, 30, or 35 at.% C showed a dark grey phase as the primary phase. The qualitative microstructure investigation showed only two phases for each sample. The XRD phase identifications showed those two phases as the metallic β(Ti,W) phase and the δ(Ti,W)C_1−x carbide. The XRD results are displayed in Figure 7; other unexpected phases could not be detected.

Samples 10 at.% C and 15 at.% C (Figure 6a,b) showed large fractions of the primary β(Ti,W) phase with a distinct signal gradient over the phase area in backscattered electron contrast. The shape of the primary phase did not show a dendritic structure. Samples 20 at.% C to 35 at.% C all showed a eutectic microstructure with either primary β(Ti,W) (Figure 6c) or primary δ(Ti,W)C_1−x (Figure 6d–f).

An exemplary EDS analysis on sample 35 at.% C showed only W and Ti in the bright phases and W, Ti, and C in the dark grey phases (

Table 4. Exemplary EDS results for both phases in sample 35 at.% C.

Phase	C at.%	W at.%	Ti at.%
Light	-	87.5 ± 1.79	12.5 ± 1.79
Dark grey	66.3 ± 0.24	12.6 ± 0.15	21.1 ± 0.19

). Combining the XRD and EDS analyses confirmed the bright phase to be the β(Ti,W) phase and the dark grey to be δ(Ti,W)C_1−x.

Through a qualitative microstructure inspection (Figure 6), it became apparent that an increase in the carbon content led to an increase in the δ(Ti,W)C_1−x carbide fraction. This was further supported by the quantitative microstructure analysis shown in the bottom part of Figure 8. The carbide fraction increased almost linearly from 26% (10 at.% C) to 79% (35 at.% C). Also presented in Figure 8 is the measured Vickers hardness for each sample. Each hardness value is the mean of 10 indentations. With increasing carbon content and carbide fraction, a significant increase in hardness was observed. The hardness increased from 681 ± 38 HV (10 at.% C) to 1898 ± 154 HV (35 at.% C). With a metallic phase fraction (β(Ti,W)) of 21% and a hardness of 1898 ± 154 HV, sample 35 at.% C was the most promising with regard to the ductile metal fraction and the hardness and was thus selected for the investigations of the W/Ti ratio influence.

3.2. Screening 2—Influence of Tungsten–Titanium Ratio

Figure 9 shows the representative microstructures of the samples from the second screening with a constant carbon content of 35 at.% and a varying W/Ti ratio. From sample 10:90 to 60:40, the Tungsten content increased (Table 2). Qualitatively, two phases were observed in all the samples. The backscattered SEM images showed a light and a dark grey phase in different occurrences. The EDS analyses showed only W and Ti for the light phase and W, Ti, and C for the dark grey phases. With increasing W content, the amount of the light phase increased, and the formation changed from individual precipitates and agglomerations at the grain boundaries to a more eutectic structure. The XRD measurements support the findings of a two-phase microstructure composed of β(Ti,W) and δ(Ti,W)C_1−x.

With increasing W content, the amount of the δ(Ti,W)C_1−x carbide decreased and the microstructure became finer. This is further supported by the quantitative microstructure analysis shown in Figure 10. Increasing the W content reduced the δ(Ti,W)C_1−x fraction from 98% (6.5 at.% W) to 74% (39 at.% W). The corresponding hardness values for the different W/Ti ratios did not show a high fluctuation. The hardness ranged from approx. 1900 HV to 2100 HV (Figure 10).

To determine the toughness of the individual compositions, the Palmqvist fracture toughness method according to ISO 28079 [14] was applied. However, due to the small sample size, the required load of >30 kgf could not be applied. To still obtain an estimate of the toughness behaviour of the compositions, a load of 0.5 kgf was used and the resulting crack formation was investigated qualitatively only as a direct comparison. Figure 11c,d also shows non-uniform crack formation at locations other than the indentation tips, which prevents a correct estimation of the Palmqvist fracture toughness [15].

Exemplary Vickers indentations for each composition are presented in Figure 11a–f. It was apparent that with an increasing W content and thus an increase in the metallic β(Ti,W) phase fraction, the cracks at the indentation tips decreased significantly. The finding here is that with an increasing W content and metallic β(Ti,W) phase, the fracture toughness increases.

4. Discussion

4.1. First Screening—Influence of Carbon Content

To investigate the influence of the carbon content on the phase formation and hardness of W-C-Ti alloys, compositions with a constant W/Ti ratio of 50:50 and varying carbon content were examined (Table 1). According to the liquidus surface projection (Figure 1), the compositions with 10–20 at.% C showed β(Ti,W) as the primary crystallization phase, whereas for samples with 25–35 at.% C, δ(Ti,W)C_1−x was supposed to be the primary phase [10]. Following primary crystallization, the residual melt solidified eutectically into β(Ti,W) and δ(Ti,W)C_1−x. The presented findings in Figure 6 and Figure 7 are in accordance with this. However, for samples 10 at.% C and 15 at.% C, there was no typical eutectic structure visible. This was assumed to be a result of β(Ti,W) solidifying directly on the already present β(Ti,W) primary crystals [16]. This finding is also supported by larger areas with different grey values of β(Ti,W) in the backscattered electron contrast (Figure 12), which indicate slight differences in chemical composition. The grey value gradients inside those larger areas are a result of the crystal segregation of the β(Ti,W) solid solution.

In addition to the primary and eutectic solidification, the microstructure investigation of the first set of compositions also showed the decomposition of the δ(Ti,W)C_1−x carbide phase. This was not present in sample 10 at.% C (Figure 12) but increased with increasing carbon content and was indicated by the lamellar structure in the dark grey carbide sections (Figure 13). According to Rudy [17], there is evidence of the precipitation of tungsten-rich metallic phases from supersaturated mono-carbide solid solutions (δ(Ti,W)C_1−x). This is also supported by Haldar et al. [11], who observed a decreasing solubility of tungsten in δ(Ti,W)C_1−x mixed carbide at lower temperatures. With decreasing temperatures, the solubility of the carbide for W decreased and thus the Ti content increased, leading to an overall harder carbide [18].

With a higher carbon content and a larger δ(Ti,W)C_1−x fraction, an increase in hardness could be achieved. It is apparent that a higher carbide fraction leads to an increase in hardness [9].

4.2. Second Screening—Influence of Tungsten–Titanium Ratio

Based on the first screening results, the composition containing 35 at.% C with a hardness of 1898 HV and a metallic phase fraction of approximately 21% was identified as the most promising candidate for further testing with varying W/Ti ratios. A hardness of around 2000 HV and a “binder” content of around 20% is in the range of commercial WC-Co hard metals and is thus a promising starting point for further experiments.

The W/Ti ratio was varied from 10:90, 20:80, 30:70, 40:60, and 50:50 to 60:40. The 50:50 sample here was identical to sample 35 at.% C in the first screening.

According to the liquidus projection presented in Figure 1, for all the investigated compositions, δ(Ti,W)C_1−x was the primary solidification phase. The residual melt subsequently solidified eutectically into β(Ti,W) and δ(Ti,W)C_1−x. The microstructure investigation presented in Figure 9a–f confirms those solidification reactions and could further be backed up by XRD measurements. Changes in the W/Ti ratio led to changes in the β(Ti,W)/δ(Ti,W)C_1−x ratio and the structure size. The carbide fraction dropped from 98% (10:90) to 74% (60:40).

Figure 14 and Figure 15 show a comparison of the 10:90 sample (Figure 14) and the 60:40 sample (Figure 15). The primary δ(Ti,W)C_1−x phase showed a significant difference in size and shape.

Despite the differences in composition and microstructure formation, the hardness values only showed a slight variation over the samples (Figure 10). This can be explained by the counteracting effects of the carbide composition and the microstructure size. With the variation in the overall W/Ti ratio, the W/Ti ratio of the carbide phase also changed, resulting in a higher W content in the low-Ti samples. Titanium carbides present a hardness of ~3200 HV [19], whereas tungsten carbides have a lower value of ~2600 HV [20]; thus, a lower Ti content in the carbide phase leads to a decrease in the hardness of the mixed carbide at a high W content and vice versa. The second effect is the microstructure refinement with a higher W content. This can be explained by varying the solidification temperatures with a varying tungsten content of the mixture. With increasing W content, the solidus temperature of the composition increased, which led to faster full solidification and thus a finer microstructure [11,21]. In addition, the higher thermal conductivity of W (164 ${\displaystyle \frac{\mathrm{W}}{\mathrm{m}\ast \mathrm{K}}}$) in comparison to Ti (11.4 ${\displaystyle \frac{\mathrm{W}}{\mathrm{m}\ast \mathrm{K}}}$) could lead to more rapid solidification and thus a finer microstructure.

For the investigated compositions, the decrease in the hardness of the carbide counteracted the increasing hardness of the finer microstructure and thus led to only a slight variation being observed in the overall hardness.

A significant difference, however, can be seen in the fracture toughness of the investigated compositions indicated by Palmqvist cracks. The samples with a higher W content showed an increase in the metallic β(Ti,W) phase as well as a more network-shaped appearance of the eutectic structure (Figure 15) that has an ability to stop crack propagation (Figure 11a–f) [9].

5. Conclusions

In this study, 11 alloy compositions of the W-C-Ti system were investigated with the aim of finding a promising composition for a hard metal that presents comparable properties to some commercial WC-Co hard metals, all while having no cobalt and a reduced tungsten content and that can be manufactured with laser-based additive manufacturing.

For this investigation, two screenings were carried out, and the samples produced were examined regarding their microstructure and hardness. In the first screening, the influence of the carbon content was investigated, and in the second screening, the influence of the W/Ti ratio on the microstructure and mechanical properties was studied.

By increasing the carbon content, the δ(Ti,W)C_1−x phase fraction and the hardness could be significantly increased (from 681 HV at 10 at.% C to 1898 HV at 35 at.% C). Sample 35 at.% C showed the most promising results in terms of its microstructure, phase fractions, and hardness.

By varying the W/Ti ratio, changes in the microstructure and in the phase fractions of β(Ti,W) and δ(Ti,W)C_1−x were observed. By increasing the tungsten content from 6.5 at.% to 39 at.%, the β(Ti,W) phase fraction increased from 2.0% to 25.7% and the microstructures became finer. The hardness values achieved for the alloys in the second screening were in the range of 1900–2100 HV as a result of microstructure refinement and the higher hardness of the Ti-rich carbide phase.

The compositions investigated in the second screening show promising results for an additively manufacturable W-C-Ti hard metal with phase fractions and properties similar to those of sintered WC-Co hard metals. However, further experiments are necessary to study the manufacturing of larger-volume samples and perform an in-depth qualification on properties such as bending strength, Young’s modulus, and fracture toughness. Other studies on a more bulk application of W-C-Ti composites produced by laser-based freeform fabrication have also reported fine eutectic structures with microhardness values of up to 1300 HV [12], while alternative sintered hard alloys such as VK8 and T15K6 [15] have achieved hardness values in the range of the samples in this study and have significantly improved tool life. These findings and the results of this study highlight the broader potential of carbide-based systems when processed with advanced manufacturing techniques.