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Article

Combined Approach to the Synthesis of WC-(Fe, Ni) Hard Alloys: Mechanical Activation and Spark Plasma Sintering

by
Gulzhaz Uazyrkhanova
1,
Yernat Kozhakhmetov
1,
Madina Aidarova
1,
Małgorzata Rutkowska-Gorczyca
2 and
Yerkezhan Tabiyeva
1,*
1
Center of Excellence “VERITAS”, D. Serikbayev East Kazakhstan Technical University, Ust-Kamenogorsk 070004, Kazakhstan
2
Faculty of Mechanical Engineering, Wroclaw University of Science and Technology, Wyb. Wyspianskiego 27, 50-370 Wroclaw, Poland
*
Author to whom correspondence should be addressed.
Crystals 2025, 15(8), 724; https://doi.org/10.3390/cryst15080724
Submission received: 22 July 2025 / Revised: 7 August 2025 / Accepted: 10 August 2025 / Published: 14 August 2025
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Abstract

This paper presents a combined approach to the synthesis of WC-(Fe, Ni) hard alloys obtained by mechanical activation and spark plasma sintering (SPS). The main attention at this stage of the work is paid to studying the evolution of the morphology and phase composition of WC-(Fe, Ni) powder mixtures during high-energy milling and their subsequent sintering by the SPS method. The study analyzed the effect of the mechanosynthesis time and the binder phase content on the change in the average particle size, the degree of defect formation, and the phase composition of the powders. It was found that an increase in the milling time to 240 min promotes the formation of the WC nanocrystalline structure and the accumulation of microdefects, which is accompanied by a decrease in the average particle size and an increase in the dislocation density. The X-ray phase analysis of the samples after SPS confirmed the preservation of the WC phase and the formation of the γ-(Fe, Ni) matrix without the formation of secondary carbide phases. The analysis of sample shrinkage showed three main stages: initial compaction, intense shrinkage, and structure stabilization. The obtained data demonstrate that optimization of the parameters of mechanical activation and SPS allow for effective control of the phase composition and morphology of WC-(Fe, Ni) powders, which opens up opportunities for their subsequent study in conditions of aggressive environments and radiation exposure.

1. Introduction

Hard alloys, or cemented carbides, are composite materials based on tungsten carbide (WC) in combination with a metallic binder phase, most often based on cobalt (Co), iron (Fe), and/or nickel (Ni) [1]. Their performance properties can be tailored by adjusting the grain size of the WC and the content of the binder component [2]. Due to their high hardness and wear resistance, hard alloys have found wide application in metal cutting, stamping, chipless machining, drilling, and stone and wood processing [3,4].
Despite active research in the direction of composition modification, WC-Co alloys remain the most common to date, with a share of over 80% on the world market [5]. However, the use of cobalt is accompanied by several serious limitations: high cost, toxicity, and carcinogenicity of WC-Co compositions [6,7,8,9,10,11], as well as radiation activity associated with the formation of long-lived 60Co isotopes under neutron irradiation, which makes their use in nuclear power impossible [12,13,14,15,16,17,18]. In addition, deterioration of properties is observed in aggressive environments due to insufficient corrosion resistance [10,11].
In this regard, an urgent task is to develop alternative cobalt-free hard alloys that retain high hardness and fracture toughness. Among the most promising candidates are metallic binders based on iron and nickel, in particular, the Fe-Ni alloy, which demonstrates an optimal balance of mechanical properties, corrosion resistance, and cost [15,16,17]. A number of recent studies have shown that WC-Fe-Ni alloys have comparable [18,19,20,21,22,23,24], and in some cases superior, strength, hardness, and resistance to aggressive effects to WC-Co [16,17,25].
At the same time, although using a low-melting metal binder allows for an increase in crack resistance and bending strength, due to the limitations of traditional synthesis technologies, there is an inevitable decrease in tungsten carbide hardness, as well as a reduction in the maximum allowable temperatures in the cutting zone, which imposes significant restrictions on cutting speeds. This circumstance necessitates attempts to develop a new group of cutting tools based on tungsten carbide [6,7,8] for metal-cutting tools intended for finishing and semi-finishing operations on ductile structural materials with an ultra-fine-grained (UFG) structure.
Analysis of the literature data indicates that one of the most effective approaches to obtaining dense nanostructured hard alloys is the spark plasma sintering (SPS) technology, which allows for a significant reduction in sintering temperature and time, avoids grain growth, and produces complex-shaped products without subsequent mechanical processing [25,26,27,28,29,30,31,32,33,34]. Unlike traditional methods, SPS speeds up the process of making materials dense by using heat from electric currents, magnetic fields, and other techniques [35,36]. Moreover, the technology has low energy consumption and high technological efficiency [37,38].
An important feature of the SPS technology is the ability to carry out sintering at very high heating rates (up to 2500 °C/min), which allows for a reduction in grain growth rate and the intensity of tungsten carbide particle decomposition [10]. The ability to directly vary the main parameters during the sintering process, which have the most significant impacts on the microstructure and phase composition of the powders (heating rate, temperature, sintering time, applied pressure, etc.), gives the SPS method greater “flexibility” in controlling the physical and mechanical properties of the resulting materials.
Despite its advantages, SPS is still limited to laboratory scales due to temperature inhomogeneity during the sintering of large-scale products. However, in recent years, the development of induction heating systems has opened up possibilities for scaling and using the method in industrial settings [39,40]. This creates prerequisites for the widespread adoption of cobalt-free WC composites obtained using SPS in the mining, machine engineering, and energy sectors.
However, it should be noted that spark plasma sintering (SPS) technology is widely used not only in the synthesis of cobalt-free hard alloys, but also in a number of other relevant areas of materials science. In particular, SPS is successfully used to obtain composite materials based on MAX phases, such as Ti3SiC2-Cu, which have increased wear resistance and improved mechanical characteristics due to the formation of intermetallic phases and carbides [41]. In addition, the method is actively used in the sintering of bioactive and biocompatible materials, including hydroxyapatite composites intended for use in implantology and tissue engineering [42]. SPS technology also demonstrates high efficiency in the formation of oxide ceramic materials, for example, based on Y2O2, CeO2, and Al2O2, with high density and a controlled nanostructure, which is especially in demand in the fields of electronics, catalysts, and functional ceramics [43]. Thus, the prospects of the SPS method are determined by its versatility, the possibility of fine-tuning the microstructure, and its scalability, which makes it a key tool in the development of modern materials for various functional purposes [43,44,45,46,47,48].
Within the framework of this work, the following is planned:
-
synthesis of cobalt-free hard alloys WC/Ni, WC/Fe, and WC/Fe-Ni using the SPS method;
-
comprehensive study of the influence of binder composition and WC grain size on the densification, microstructure, and hardness of the obtained composites.
The novelty of the research lies in the comprehensive evaluation of the mechanical properties and microstructure of WC composites obtained by the SPS method with alternative binders.

2. Materials and Methods

The study was conducted on micron-sized tungsten carbide (WC) powders produced by Hebei Suoyi New Material Technology Co., Ltd., Handan, China. According to the manufacturer’s specifications, the initial particle sizes were in the range of 30–50 μm. The iron and nickel nanopowders were obtained using the plasma synthesis method in a direct current arc with reduction annealing in a hydrogen atmosphere. The choice of micron-sized tungsten carbide powders was driven by the goal of obtaining fine-grained hard alloys with improved mechanical properties.
The general chemical composition of the powders under study is presented in Table 1. The chemical composition of the factory WC and Fe, Ni powders is given according to the supplier’s certificates (Hebei Province).
Six samples with different chemical compositions of alloying elements are compared.
A series of experiments on the mechanosynthesis of a powder system based on WC with various alloying components (Fe, Ni) has been carried out. Table 2 displays the conditions of the conducted mechanosynthesis experiments.
The loading of components into the grinding jar was carried out in a VGB-3C “Ar atmosphere glove box”. The mechanosynthesis of powder mixtures of the WC system with alloying components was carried out on a Pulverisette 7 Premium line planetary ball mill, Fritsch GmbH, Idar-Oberstein, Germany, where the main parameters were acceleration, duration, and ball-to-powder ratio.
Mechanosynthesis resulted in the formation of six powder mixtures based on WC. The obtained mixtures were sorted into separate flasks and labeled with an indication of the processing parameters. These flasks with powder mixtures were sent for subsequent study by SEM and X-ray fluorescence methods; then these powder mixtures were sintered on the SPS unit.
For the mechanical activation process, precisely weighed WC, Fe, and Ni powders were added to grinding cups with stainless steel grinding balls, which were placed in a planetary micro-mill. The ratio of the grinding balls in the powder composition was 10:1 at a rotational speed of 500.
During high-energy processing, stearic acid was added to all powder mixtures in an amount of 3% of the powder weight, acting as a surfactant and a process agent. The additive in the specified concentration helps optimize the mechanical alloying process by reducing the agglomeration of particles, decreasing friction between them, and increasing the homogeneity of mixing. At the same time, the content of stearic acid at the level of 3% is optimal, since it ensures an improvement in the processability of the process without a significant effect on the chemical composition and phase composition of the metal powder mixture, which is confirmed by the results of X-ray phase and elemental analysis [49].
The changes in the shape and structure of the samples were examined with a scanning electron microscope (SEM, Tescan Mira, TESCAN GROUP, a.s., Brno, Czech Republic) in a mode that looks at both composition and surface details. The size of the particles was measured with a laser diffraction analyzer (Analysette-22 NeXT Nano, Fritsch, Fritsch GmbH, Idar-Oberstein, Germany), and the types of phases and structural features of the powders were analyzed using an X-ray diffractometer X PertPRO (Malvern Panalytical Ltd., Almelo, The Netherlands) with an anode-Cu, measurement angle of 20–100°, X-ray tube voltage of 45 kV, current of 40 mA, step of 0.02, and measurement time of 1.5 s at each step. The crystallite size was determined using mathematical calculations based on the Scherrer equation:
D = K λ β c o s θ
where D is the size of the crystallite, K is the shape coefficient, λ is the wavelength of the X-ray radiation, β is the total peak width at half the height (FWHM), and θ is the diffraction angle. The peak widths were determined by fitting the Gaussian function in the High Score program.
In the calculations, a shape factor (K-factor) of 0.9 was used, which is a standard value for spherical or nearly spherical crystallites when analyzed in reflected geometry. No correction for instrumental broadening was made, since its contribution, according to the technical characteristics of the X’Pert PRO X-ray diffractometer (Malvern Panalytical Ltd., Almelo, The Netherlands), is less than 0.1°, which is significantly lower than the total broadening of the observed diffraction maxima (in the range of 0.3–0.6°) in the samples under study. Thus, the effect of instrumental broadening on the calculated values of crystallite sizes and microstrains was recognized as insignificant.
The sintering process of powder mixtures was carried out on the Dr. Sinter SPS-625 installation in a graphite mold at a temperature of 1200 °C, a pressure of 70 MPa, and a heating rate of 50 °C/min in vacuum (2−5 Pa) without isothermal exposure. The temperature was recorded by a radiation pyrometer, and shrinkage and its rate were recorded by an integrated dilatometer.
This work used the hydrostatic weighing method to determine the density and porosity of the samples. Hydrostatic weighing is a method for measuring the mass per unit volume of an object being studied.
In this work, high-precision Shimadzu AUW220D (Shimadzu Corporation, Kyoto, Japan) analytical scales were used to determine the density and porosity of the samples. The measurement procedure includes several steps. First, the mass of the dry samples was determined, after which the samples were immersed for 24 h in a container with distilled water. The next step was to remove the test samples from the water and place them on an analytical scale, while measuring the mass of samples saturated with water. At the final stage, the mass of the studied samples was measured when they were completely immersed in a tank filled with distilled water.

3. Results and Discussion

Figure 1 and Figure 2 show the X-ray diffraction spectra of WC-Ni-Fe powders with mass ratios of 88:6:6 and 85:7.5:7.5 subjected to grinding in a planetary ball mill at 500 rpm for various times (0, 1, 2, and 4 h). At the start (0 h), both mixtures show strong and sharp diffraction peaks that match the hexagonal phase WC, along with the phases of metallic tungsten (W), iron (Fe), nickel (Ni), and W2C carbide. This pattern indicates a high degree of crystallinity and the presence of a multicomponent structure.
After one hour of grinding, both compositions show a noticeable broadening of their diffraction peaks and a decrease in their intensity, especially for the WC phase, which indicates a decrease in the average size of crystallites, an increase in microdeformations, and an increase in the density of lattice defects. The formation of new phases has not been recorded, which confirms the predominantly physical nature of the transformations taking place.
Grinding for 2 to 4 h leads to a significant drop in intensity and a further widening of the WC peaks, showing that the material is moving towards a nanocrystalline state and that more defects are building up in the crystal structure. The mechanisms underlying these changes include the development of plastic deformation, the formation of sub-boundaries, an increase in dislocation density, and possible surface amorphization.
The calculated values of crystallite sizes and microdeformations of WC88Ni6Fe6 powders confirm the above observations. The crystallite size decreases from 517 Å (0 h) to 265 Å (4 h), and the microdeformation of the crystal lattice increases from 0.185% to 0.35% (Table 3). At the same time, for WC85Ni7.5Fe7.5 powders, the crystallite size decreases from 489 Å (0 h) to 222 Å (4 h), and the microdeformation of the crystal lattice increases from 0.196% to 0.36% (Table 3). Both powder compositions are characterized by a steady decrease in the average size of crystallites with a simultaneous increase in microdeformations as the grinding time increases.
These changes in structural parameters are consistent with the results of the X-ray diffraction analysis and reflect the accumulation of defects in the crystal lattice caused by plastic deformation during high-energy grinding. The data from the detailed analysis of the structural parameters show that the powders studied are changing into a nanocrystalline state and that significant changes in their structure occur as the processing time increases.
A comparative analysis showed that WC85Ni7.5Fe7.5 powders are characterized by a higher grinding rate and a higher intensity of defect formation processes compared to WC88Ni6Fe6. Already at the initial stage of grinding, lower values for the average size of crystallites and higher values for microdeformations are recorded for this composition, which is probably due to the increased content of plastic components in the bundle, contributing to a more efficient transfer of mechanical energy during collisions.
The analysis of the morphology and particle sizes (see Figure 2) of the WC-Ni-Fe powder mixtures, carried out using scanning electron microscopy (SEM) and statistical data processing, revealed characteristic differences in the dynamics of grinding between the two studied compositions—WC85Ni7.5Fe7.5 and WC88Ni6Fe6—during processing in a planetary ball mill with a rotation speed of 500 rpm. At the initial stages of mechanical impact (20 min), both compositions show a significant reduction in the average particle size compared to the initial state (see Table 4), which reflects the high efficiency of the processes of agglomerate fragmentation and initial deformation. The average particle size at this stage is ~9.6–9.8 microns for both systems.
However, with an increase in the treatment time to 60 and 120 min, an unexpected increase in the average particle size is observed, which can be attributed to the effect of secondary aggregation. The formation of aggregates results from both the plastic deformation of the binding phase and the enlargement of particles caused by the adhesion of finely dispersed fragments.
The observed increase in the average particle size with increasing grinding time to 60 min, and then its repeated decrease, can be explained by processes of dynamic reagglomeration. Similar fluctuations in particle size at advanced stages of mechanical processing were previously recorded and described, including in the work of Dvornik and Zaitsev [38], where the effect of intense plastic deformation on the formation of agglomerates with a metastable nature is emphasized.
This nonlinear relationship between milling time and particle size is in good agreement with the literature data, which reports the alternation of competing grinding and agglomeration processes during high-energy processing [49].
The most significant differences between the systems appear during long-term processing (240 min). For WC85Fe7.5Ni7.5, the average particle size decreases to 2.71 ± 0.03 microns, while for WC88Fe6Ni6 it is 3.96 ± 0.26 microns. This variation indicates a more efficient destruction of particles and the destruction of agglomerates in an alloy with a high content of nickel and iron. The SEM images (see Figure 3) support this finding: after 4 h of processing, the particles in WC85Fe7.5Ni7.5 are more broken apart, less round, and have a more even size, while WC88Fe6Ni6 still has some larger and only slightly broken clusters.
Presumably, this difference is due to the more pronounced plasticity of the binding phase in WC85Fe7.5Ni7.5, which contributes to a more intense transfer of impact energy to tungsten carbide particles and, as a result, accelerated particle destruction. According to [50,51], an increase in the proportion of Ni and Fe in the bond helps reduce the hardness and increases the deformability of the bonding phase, which optimizes the grinding process.
So, the mixture WC85Fe7.5Ni7.5 shows better stability when under long-term mechanical stress, leading to smaller and more evenly shaped particles. This makes it preferable in the development of composite materials where high packing density and minimal porosity are critically important, for example, during subsequent spark plasma sintering.
All six WC–Ni–Fe powder mixtures, differing in composition (WC85Fe7.5Ni7.5 and WC88Fe6Ni6) and pre-treatment time in a planetary ball mill (60, 120, and 240 min), were subjected to SPS.
All samples demonstrated a classic three-stage shrinkage curve (see Figure 4): I—slight compaction (600–900 °C), II—intense shrinkage (1000–1200 °C), and III—stabilization when approaching the maximum temperature. For powders treated for 60 min, the start of stage II was recorded at ~1040 °C, the maximum shrinkage rate was 6.2·10−3 mm/s, and the final linear shrinkage was about 14.8%. At 120 min of grinding, the beginning of intensive shrinkage shifted to ~1000 °C, the maximum speed increased to 7.5·10−3 mm/s, and shrinkage reached 16.5%. The most pronounced effect was observed after 240 min, where shrinkage began already at ~960 °C, reaching 9.2·10−3 mm/s, and the total deformation was 18.3%. This may be attributed to the high reactivity of the powders and the increased surface area resulting from a decrease in particle size during processing.
X-ray phase analysis (XRD) performed after the spark plasma sintering of all six samples showed the presence of a stable WC carbide phase in the form of a hexagonal structure with slight peak shifts (see Figure 5), which may indicate mechanical alloying and residual stresses after high-energy treatment. The bonding phase was predominantly a γ-phase with a cubic lattice of the type γ-(Fe,Ni), forming a solid metal matrix between the WC grains.
For WC85Fe7.5Ni7.5 compositions, weak diffraction peaks corresponding to the intermetallic compound Fe3W3C (η phase) were recorded in the samples after 60 min of mechanical activation, which is consistent with the data [51,52,53], where a tendency to form η-phases is noted in WC–Fe(Ni) systems with an excess of iron and a lack of free carbon. Increasing the milling time to 120 and 240 min eliminates signs of the presence of the η-phase, probably due to improved homogenization of the composition and uniform distribution of carbon in the mixture.
The WC88Fe6Ni6 compositions showed no signs of Fe3W3C formation, indicating a more stable phase equilibrium at the given Fe/Ni ratio. This result confirms the conclusions of [53], in which the absence of the η-phase in similar compositions is explained by the reduced width of the “carbon window” in the WC–Fe–Ni systems and better stabilization of the γ-phase with increased nickel content.
The disappearance of the W2C peaks after sintering indicates a phase transition in which W2C reacts with the Ni–Fe solid solution, probably forming a more stable phase at elevated temperatures. This process is driven by the diffusion of atoms across particle boundaries at temperatures below the melting point, a fundamental aspect of sintering in which particles fuse together. The reaction forms a new solid phase, absorbing the original W2C and causing it to disappear from the diffraction pattern.
Thus, phase analysis confirms the following:
-
stability of WC as the main solid phase;
-
formation of γ-phase plastic (Fe, Ni) as an effective binder;
-
elimination of η-phase in optimally ground and sintered systems;
-
absence of secondary carbides with the correct choice of Fe/Ni ratio and sintering mode.
Based on the analysis of the results of the microstructure and phase composition of the sintered samples, a summary analysis of the phase composition of the sintered samples was formed, reflecting the trends in the presence of the η-phase and the homogeneity of the phase structure (see in Table 5):
The density of all samples obtained after SPS was measured by hydrostatic weighing using a standard measuring complex (ISO 3369:2006 [54]). The results are presented in Table 6.
Data analysis shows that increasing the grinding time has a positive effect on the density of the samples for both compositions. The WC88Fe6Ni6 sample achieves the highest relative density (98.6%) after 240 min of pretreatment. This increase is due to improved intergranular compaction and a more uniform distribution of the Ni-Fe binder phase. For WC85Fe7.5Ni7.5, the density also increases with longer grinding times, but the maximum value reaches only 97.9% due to more pronounced agglomeration of the WC grains and lower fluidity of the binder component.
Although this paper focuses on the microstructure and phase composition of sintered samples, measurements of Vickers hardness, flexural strength, and fracture toughness are currently being actively conducted. These data, including fracture toughness evaluation results, will be presented in a future publication.
Obtaining a full range of mechanical characteristics, including hardness, bending strength, crack resistance, and wear resistance, will allow us to approach the solution to the problem of creating a highly efficient ceramic cutting tool designed to operate in a wide range of thermal and force cutting conditions.
Currently, certain research teams are initiating efforts to develop a metal-cutting tool (MCT) utilizing the SPS approach. The absence of models for high-speed sintering of tungsten carbide, along with models elucidating the correlations between the sintering parameters, microstructure, and properties of these ceramics, hinders researchers from attaining significant advancements in this field and progressing towards the development of an industrial technology for the production of highly efficient MCT.

4. Conclusions

During the study of WC–(Fe, Ni) powder mixtures obtained by mechanical activation, it was hypothesized that increasing the milling time to 240 min may contribute to a reduction in the average particle size and the development of microdefects within the crystalline structure. This process appears to be accompanied by the potential formation of a nanocrystalline structure and a possible increase in dislocation density, as suggested by the results of X-ray phase analysis.
The analysis of the phase composition evolution of WC–(Fe, Ni) powders during mechanosynthesis suggests that the hexagonal WC phase is likely preserved, while a γ-(Fe, Ni) matrix may form. This matrix is assumed to support a more uniform distribution of the metallic binder among tungsten carbide grains, potentially without the formation of secondary carbide phases.
Preliminary spark plasma sintering experiments indicate a three-stage shrinkage behavior in all samples: initial compaction (600–900 °C), intensive shrinkage (1000–1200 °C), and structural stabilization near the maximum temperature (1200 °C). An increase in mechanosynthesis time to 240 min appears to correlate with a higher shrinkage rate of up to 9.2·10−3 mm/s, which may be attributed to a decrease in particle size and an increase in powder reactivity.
X-ray phase analysis of the sintered samples tentatively confirms the preservation of the WC phase and the formation of a γ-(Fe, Ni) matrix. The absence of the η-phase Fe3W3C under the applied sintering conditions suggests that these parameters may be suitable for the studied system.

Author Contributions

Conceptualization, G.U., Y.K., M.A. and Y.T.; Formal analysis, Y.T. and M.A.; Methodology, Y.K., M.R.-G., G.U., M.A. and Y.T.; Supervision, Y.K.; Visualization, Y.K., M.R.-G. and G.U.; Writing—original draft, Y.K., G.U., M.R.-G., M.A. and Y.T. All authors have read and agreed to the published version of the manuscript.

Funding

This research is funded by the Ministry of Science and Higher Education of the Republic of Kazakhstan. (Grant No. BR24992925).

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author(s).

Conflicts of Interest

The authors declare no conflicts of interest.

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Crystals 15 00724 g001aCrystals 15 00724 g001b
Figure 1. Evolution of the phase composition of powder mixtures at 500 rpm depending on the processing time: (a) WC88Ni6Fe6; (b) WC85Ni7.5Fe7.5.
Figure 1. Evolution of the phase composition of powder mixtures at 500 rpm depending on the processing time: (a) WC88Ni6Fe6; (b) WC85Ni7.5Fe7.5.
Crystals 15 00724 g001aCrystals 15 00724 g001b
Crystals 15 00724 g002aCrystals 15 00724 g002b
Figure 2. Particle size distribution chart of samples.
Figure 2. Particle size distribution chart of samples.
Crystals 15 00724 g002aCrystals 15 00724 g002b
Crystals 15 00724 g003
Figure 3. The SEM images of powder mixtures at different processing times.
Figure 3. The SEM images of powder mixtures at different processing times.
Crystals 15 00724 g003
Crystals 15 00724 g004
Figure 4. Characteristic image of shrinkage of WC85Fe7.5Ni7.5 alloy samples during sintering.
Figure 4. Characteristic image of shrinkage of WC85Fe7.5Ni7.5 alloy samples during sintering.
Crystals 15 00724 g004
Crystals 15 00724 g005
Figure 5. The phase composition of the samples after sintering: (a) WC88Ni6Fe6; (b) WC85Ni7.5Fe7.5.
Figure 5. The phase composition of the samples after sintering: (a) WC88Ni6Fe6; (b) WC85Ni7.5Fe7.5.
Crystals 15 00724 g005
Table 1. Chemical composition of the initial powders.
Table 1. Chemical composition of the initial powders.
Chemical Composition of Powders, at.%
FeCrBiAlZnMnCuCoCaNiCSO
0.0020.0010.0010.0020.0020.0010.0030.0010.0020.0020.0010.0005
NiSiFe Zr
0.0030.0010.001 0.0020.0010.025 0.002
WC B F
0.0010.0030.0020.0010.0010.0030.0040.0020.0020.0030.001≤0.040
Table 2. Modes for mechanical activation of WC-FeNi, WC-Fe, and WC-Ni and chemical composition of the powder mixture.
Table 2. Modes for mechanical activation of WC-FeNi, WC-Fe, and WC-Ni and chemical composition of the powder mixture.
Sample NameTime, minRotation Speed, rpmEnvironment
Mixing powder mixtures20200Argon
Mechanosynthesis
Sample 1
88WC-6Fe-6Ni at.%		60500Argon
120
240
Sample 2
85WC-7.5Fe-7.5Ni at.%		60
120
240
Table 3. Microdeformation of the crystal lattice of powder mixtures WC88Ni6Fe6 and WC85Ni7.5Fe7.5 at 500 rpm.
Table 3. Microdeformation of the crystal lattice of powder mixtures WC88Ni6Fe6 and WC85Ni7.5Fe7.5 at 500 rpm.
WC88Ni6Fe6WC85Ni7.5Fe7.5
Milling timeLattice strain [%]Crystallite size [Å]Milling timeLattice strain [%]Crystallite size [Å]
0 h0.1855170 h0.196489
1 h0.282951 h0.277282
2 h0.3252802 h0.318253
4 h0.352654 h0.358222
Table 4. Arithmetic mean of particle size of powder mixtures.
Table 4. Arithmetic mean of particle size of powder mixtures.
Duration of the ProcessWC88Ni6Fe6WC85Ni7.5Fe7.5
initial27.25 ± 5.66 microns
20 min9.5 +/− 0.45 microns9.82 +/− 0.11 microns
60 min13.7 +/− 0.53 microns13.42 +/− 0.18 microns
120 min12.57 +/− 0.39 microns12.02 +/− 0.22 microns
240 min3.96 +/− 0.25 microns2.71 +/− 0.03 microns
Table 5. Phase composition of sintered samples.
Table 5. Phase composition of sintered samples.
SampleGrinding Time (min)Main PhaseMetal MatrixPresence of η-Phase (Fe3W3C)Homogeneity of Phases
WC85Fe7.5Ni7.560WCγ-(Fe,Ni)Presencelow
WC85Fe7.5Ni7.5120WCγ-(Fe,Ni)Traces of the presenceaverage
WC85Fe7.5Ni7.5240WCγ-(Fe,Ni)No η-phasehigh
WC88Fe6Ni660WCγ-(Fe,Ni)No η-phaseaverage
WC88Fe6Ni6120WCγ-(Fe,Ni)No η-phasehigh
WC88Fe6Ni6240WCγ-(Fe,Ni)No η-phasehigh
Table 6. The density of samples obtained by SPS.
Table 6. The density of samples obtained by SPS.
SampleGrinding Time (min)Theoretical Density, g/cm3Measured Density, g/cm3Relative Density, %
WC85Fe7.5Ni7.56014.513.895.2
WC85Fe7.5Ni7.512014.514.096.6
WC85Fe7.5Ni7.524014.514.297.9
WC88Fe6Ni66014.714.195.9
WC88Fe6Ni612014.714.397.3
WC88Fe6Ni624014.714.598.6
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Uazyrkhanova, G.; Kozhakhmetov, Y.; Aidarova, M.; Rutkowska-Gorczyca, M.; Tabiyeva, Y. Combined Approach to the Synthesis of WC-(Fe, Ni) Hard Alloys: Mechanical Activation and Spark Plasma Sintering. Crystals 2025, 15, 724. https://doi.org/10.3390/cryst15080724

AMA Style

Uazyrkhanova G, Kozhakhmetov Y, Aidarova M, Rutkowska-Gorczyca M, Tabiyeva Y. Combined Approach to the Synthesis of WC-(Fe, Ni) Hard Alloys: Mechanical Activation and Spark Plasma Sintering. Crystals. 2025; 15(8):724. https://doi.org/10.3390/cryst15080724

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Uazyrkhanova, Gulzhaz, Yernat Kozhakhmetov, Madina Aidarova, Małgorzata Rutkowska-Gorczyca, and Yerkezhan Tabiyeva. 2025. "Combined Approach to the Synthesis of WC-(Fe, Ni) Hard Alloys: Mechanical Activation and Spark Plasma Sintering" Crystals 15, no. 8: 724. https://doi.org/10.3390/cryst15080724

APA Style

Uazyrkhanova, G., Kozhakhmetov, Y., Aidarova, M., Rutkowska-Gorczyca, M., & Tabiyeva, Y. (2025). Combined Approach to the Synthesis of WC-(Fe, Ni) Hard Alloys: Mechanical Activation and Spark Plasma Sintering. Crystals, 15(8), 724. https://doi.org/10.3390/cryst15080724

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