# Broadening of the carbon window and the appearance of core-rim carbides in WC-Fe/Ni cemented carbides

Victor Lamelas, Manon Bonvalet Rolland, Lisa Toller-Nordström, Raquel de Oro Calderón, Martin Walbrühl, Annika Borgenstam

► To cite this version:

Victor Lamelas, Manon Bonvalet Rolland, Lisa Toller-Nordström, Raquel de Oro Calderón, Martin Walbrühl, et al.. Broadening of the carbon window and the appearance of core-rim carbides in WC-Fe/Ni cemented carbides. Journal of Alloys and Compounds, 2024, Journal of Alloys and Compounds, 999, pp.175078. 10.1016/j.jallcom.2024.175078. hal-04991707

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# Broadening of the carbon window and the appearance of core-rim carbides in WC-Fe/Ni cemented carbides

V. Lamelas<sup>a, *</sup>, M. Bonvalet Rolland<sup>a, b</sup>, L. Toller-Nordström<sup>a</sup>, R. de Oro Calderón<sup>c</sup>, M. Walbrühl<sup>d</sup>, A. Borgenstam<sup>a</sup>

<figure><img src="image_3.png" /></figure>

## ARTICLE INFO

### Keywords:

ICME

Cemented carbides

$\eta_6$-carbide

$\kappa$-carbide

Alternative binder

Modelling

## ABSTRACT

Among several separate challenges, the major one for replacing cobalt in cemented carbides is the difficulty to obtain alternative binder materials with a C-window broad enough to be robustly processed under conventional industrial control on the C content. The C-window is defined as the C content range for which phases that are detrimental to the mechanical properties are avoided. The present paper has two main objectives: first, to show that the processing C-window of Fe-Ni based systems is in fact wider than what thermodynamic equilibrium calculations predict, and that its width can be controlled moderately by tweaking the initial WC grain size and the cooling rate used in the material's processing. Secondly, in case those detrimental phases are not avoided, this work gives insight on how to make their appearance less detrimental for the mechanical properties. The morphology, volume fraction and particle size distribution of the detrimental phases, specifically $\eta_6$-carbides at low C contents, are investigated to explore desirable combination of hardness and toughness of alternative binder cemented carbides. During this study it was also discovered that in samples with carbon contents below the low-C limit of the C window a carbide with hexagonal lattice known as $\kappa$, not commonly seen in cemented carbides, appeared and formed the core of a core-rim structure together with the more common $\eta_6$-phase. It is believed that the $\kappa$-carbide form due to local high concentrations of tungsten during solid state sintering and that it has an impact on the precipitation characteristics of the $\eta_6$-phase.

## 1. Introduction

Cemented carbides are composite materials that combine the high hardness of carbide reinforcing particles, usually tungsten carbide, WC, with the toughness of a metal binder, usually cobalt, Co. The optimal combination of both properties makes them the perfect choice for many industrial applications where wear resistance and heavy-duty performance are required. The key role that cemented carbides have in a changing industrial environment characterized by the continuous search for higher efficiency, better performance, and greener solutions, requires the development of new material grades to meet these expectations [1]. For instance, the need of replacing Co from cemented carbides has been a vastly researched topic in the last decades [2,3,4,5,6], driven by the potential changes in toxicity categorization of Co powders [7] as well as the need to reduce the industry dependence on critical raw

materials [8]. However, Co-based cemented carbides present several advantages when compared to cemented carbides with alternative binder alloys that are challenging to match. Notably, Co exhibits superior wettability compared to its Co-free counterparts, which enhances its sinterability. Co as a binder also displays excellent mechanical properties due to the combination of a significant dissolution of W during sintering (benefiting the solution hardening) [9] as well as the formation of dense interwoven stacking fault networks resulting from the metastability of FCC-Co at room temperature [10]. Further, the main obstacle for the replacement of Co as the primary component of cemented carbides binder has been its comparatively broader C-window with respect to alternative binder systems. The C-window is defined as the C range where the formation of phases that are detrimental for mechanical properties are avoided. In this sense, cemented carbides having a broad C-window are beneficial since the obtained microstructure and therefore

<footnote><sup>a</sup> KTH Royal Institute of Technology, Department of Materials Science and Engineering, Brinellvägen 23, Stockholm SE-100 44, Sweden</footnote>

<footnote><sup>b</sup> Univ. Lille, CNRS, INRAE, Centrale Lille, UMR 8207 - UMET - Unité Matériaux et Transformations, Lille F-59000, France</footnote>

<footnote><sup>c</sup> TU Wien, Vienna 1060, Austria</footnote>

<footnote><sup>d</sup> AB Sandvik Coromant R&D, Lerkrogsvägen 19, Stockholm SE-126 80, Sweden</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: vlc@kth.se (V. Lamelas).</footnote>
its properties will be robust to eventual compositional fluctuations inherent to the manufacturing process, in the order of $\Delta \le 0.05$ wt%C [11].

Given the significant role of the C-window in determining if a cemented carbide grade can be industrially processed and thus be a feasible replacement of Co, numerous studies have investigated the C-window of various alternative binder systems [12,13,14,15,16]. In the search for materials with a suitable C-window ($\Delta>0.05$ wt%C), computational methodologies like Integrated Computational Materials Engineering (ICME) have aided researchers to explore new compositional landscapes at a faster pace. Within this framework, the CALPHAD methodology [17] has been used to calculate phase diagrams that can be used as first approximation to define the C-window as a C content range where only FCC and WC are in equilibrium. According to this definition, the C-window is usually defined in the temperature vs wt%C isopleth of the relevant system as the C range between the WC+Liq+$\eta_6$+FCC and WC+Liq+Graph.+FCC invariant regions, depicted as red crosses in Fig. 1. The definition of such limits has greatly assisted the development of new grades. However relying only on equilibrium data to define the C-window introduces some uncertainties that are worth studying in further detail.

The limits of the equilibrium C-window depend on the temperature at which kinetic aspects during the cooling from sintering temperature are considered relevant. When only using information from the phase diagram, different C-windows can be defined depending on if only liquid-solid phase transformations are considered (red crosses in Fig. 1) or if the mobility of elements is considered relevant down to a freeze-in temperature around 1000 °C (blue crosses in Fig. 1). The need to accurately predict such temperature limits favors the definition of a dynamic carbon window that uses information on the material's kinetics in addition to equilibrium predictions to assess the impact of kinetics on the formation of third phases. Even in well-established systems like WC-Co, a considerable impact on the width of the C-window has been reported by the present authors when considering experimental processing conditions [18]. The C content range where a two-phase material is obtained can be nearly twice the range predicted by thermodynamic equilibrium calculations when coarse WC grains and fast cooling rates are employed in the liquid phase sintering process. Understanding if similar results can be achieved when processing alternative binder alloys, with narrower C-windows according to equilibrium, could determine whether a binder system is categorized as “industrially

unprocessable” or “potential candidate for Co replacement”. This holds particular significance in compositions with carbide-forming elements like Fe, where the roofing of the phase diagram is promoted (point L in Fig. 1, moving towards wt%C > point H), and the narrowing of the equilibrium C-window is accentuated.

In the current work, an ICME-based modeling approach is applied to understand how processing kinetics affect the C-window limits in cemented carbides with Fe-Ni based binders. Selected experimental characterization on WC + 11 wt% Fe-Ni 1:1 (weight ratio 1–1) cemented carbide is also presented as a key factor for model calibration and identification of phase transformation mechanisms. The selection of such binder composition is based on both technological and practical motivations. On one hand, Fe-Ni 1:1 binder retains the grain growth inhibition (GGI) effects inherent to Fe based binders [19,20] and it keeps an FCC crystallography at room temperatures (both advantageous for mechanical properties). On the other hand, the slightly wider equilibrium C-window compared to pure Fe binders, simplifies the processing of these cemented carbides under experimental conditions. Furthermore, parameters used in precipitation modeling, such as interfacial energies or wetting angles, vary with the nominal composition. Therefore, calibrating the model using this intermediate Fe-Ni ratio is deemed an efficient approach that saves additional experimentation.

Special focus is given to the low C limit in the modeling of the Dynamic C-window of Fe-Ni systems. First, because at low carbon contents, the grain growth of WC grains has been shown to be minimized [21,22], making this end of the C-window of higher interest. Secondly, the effect of processing kinetics on lowering the low-C end has also been proven to be much stronger than on moving the C-limit bounded by graphite precipitation in Co-based cemented carbides [18]. Finally, at low C contents controlling the formation of $\eta_6$-carbides under adequate processing conditions has potential to act as beneficial reinforcing precipitates in the microstructure.

The presence of $\eta_6$ is for instance well-known to enhance the material properties in cryogenic treated cemented carbides as it precipitates as finely distributed hard-particles [23,24,25,26]. Cryogenic treated cemented carbides tend to report higher hardness and comparable toughness compared to untreated materials, leading to an enhanced wear-resistance. The achievement of microstructures as fine as those obtained in cryogenic-treated samples directly after sintering conditions seems challenging to obtain due to the high temperatures in the processing of cemented carbides that promote coarser microstructures.

<figure><img src="image_4.png" /><figcaption>**Fig. 1.** Isopleth of a) WC + 1:1 Fe-Ni with 11 wt% binder. b) WC + 11 wt% Co. The limits of the C-window for each material are defined by the red crosses. Stars correspond to the experimentally tested C contents in the present study.</figcaption></figure>
However, such studies show that the detrimental impact of η₆-carbides on cemented carbides mechanical properties can be considered a function of their size and distribution. Samples with smaller and well distributed η₆-carbides are better than those with dense concentrations [5]. A deeper understanding of the formation mechanisms of η₆-carbides and the relationship that their tailored microstructure has over the cemented carbides mechanical properties could be used to enhance their performance, as well as to broaden the range of binder systems that previously showed more restrictive C-window definitions. In opposition to the traditionally defined C-window, the definition of a dynamic C-window does not necessarily aim to avoid the precipitation of third phases to attain a two-phase microstructure (WC+FCC) by default, instead searches for the adequate set of processing conditions (composition, cooling rate, sintering temperature, grain size etc.) that lead to a microstructure giving suitable properties.

In this work, the Dynamic C-window model presented in previous work [18] is applied to Fe-Ni based systems. Experimental characterization results on Fe-Ni 1:1 samples iterating on the process cooling rate, initial grain size and sintering temperature are shown for model calibration purposes, as well as to map the diversity of microstructures that can be obtained and to further study in detail the possible formation mechanisms of η₆-carbides. Following an ICME methodology, the effects of processing conditions on the precipitation limits observed in Fe-Ni 1:1 system are extrapolated with reasonable reliability for the whole Fe-Ni binder-based system.

## 2. Materials and methods

### 2.1. Modeling of η₆-carbide precipitation in alternative binder cemented carbides

The microstructural evolution of cemented carbides during the cooling step of the sintering process is described using the thermo-kinetic modelling scheme presented in [18,21]. The coarsening of WC and the formation and possible coarsening of η₆-carbides as temperature decreases from the sintering temperature is modeled, first in a liquid matrix and second, after describing the solidification process, in a solid matrix. The main contribution of the proposed strategy relies on the interactive coupling between several models and tools to describe and understand the obtained microstructure at room temperature. The Thermo-Calc (TC) software [27] is used for all thermodynamic calculations. The data extracted from thermodynamic and kinetic databases (TCFE10 and MOBF5 [28,29]) is then used as an input for a mean-field approach model (based on [30,31]) that tracks the formation/dissolution of detrimental phases (η₆) and the coarsening of WC grains. In the latter, the size distribution of each of these phases is modeled describing the nucleation, growth and coarsening of precipitate classes within the Langer-Schwartz-Kampmann-Wagner approach (LSKW). In addition to these tools, the Scheil-Gulliver (S-G) tool in Thermo-Calc [32] is used to describe the solidification process below the peritectic temperature. Solidification results serve as an input for further mean-field calculations in the solid matrix. The reader is addressed to reference [18] for more details.

This modelling framework has been adapted to include some peculiarities found when processing Fe-Ni based cemented carbides. Among the peculiarities, the modeling tool has been adapted to account for the presence of third phases at the end of the sintering holding time, which marks the onset of the cooling and serves as the initial state in the simulations. Additionally, the model now also accounts for potential segregation profiles in η₆-carbides, formed along cooling, to calculate η₆ composition.

The parameters used for the simulations are summarized in Table 1, where σ<sub>Matr.-Precip.</sub> and θ<sub>Precip.-Substrate</sub> are the interfacial energy between the matrix and the precipitate and the dihedral angle between the precipitate and WC particles, respectively. The dependence of the geometry on the nucleation barrier resulting from Fe-Ni 1:1 system's

<table><caption>**Table 1**

Parametrization used for the modeling of η₆ precipitation, obtained from the adjustment to experimental results.</caption><thead><tr><th>Temperature range</th><th>T<sub>Sint</sub> → T<sub>Solid</sub></th><th></th><th>T<sub>Solid</sub> → T<sub>Freeze-in</sub></th></tr></thead><tbody><tr><td>Matrix</td><td>Liquid</td><td></td><td>FCC</td></tr><tr><td>Precipitates</td><td>η<sub>6</sub> (M<sub>6</sub>C)</td><td>WC</td><td>η<sub>6</sub> (M<sub>6</sub>C)</td></tr><tr><td>σ<sub>Matr.-Precip.</sub> (J/m<sup>2</sup>)</td><td>1.35</td><td>0.14</td><td>TC-calculated<br>f(T) ≈ 0.17 – 0.18</td></tr><tr><td>θ<sub>Precip.-Substrate</sub> (°)</td><td>14</td><td>-</td><td>45</td></tr></tbody></table>

parameterization showed a close resemblance to the one obtained for Co-based systems [18].

### 2.2. Sample processing

All the investigated alloys contain WC with 11 wt% Fe-Ni binder with a weight ratio of 1–1 and can be positioned in the isopleth shown in Fig. 1.a. The nominal C content was controlled through the addition of graphite or pure tungsten powders accordingly. The target C contents (legend in Fig. 2) were compensated according to the usual C-loss calibrated for the used furnace, specifically 0.1 wt% C was added to the powder blends. Two experimental matrixes with different aims have been processed:

The first, being the most complete experimental matrix, aims to set the compositional limits (carbon window) for which a two-phase cemented carbide is obtained as a function of the grain size and cooling rate. To determine this, a set of C contents ranging from the η-phase to the graphite containing region, with a higher number of samples for the low-carbon contents was chosen.

The effect of the grain size is studied through the processing of two sets of samples with a ranging C content, with two well differentiated mean sizes. From now on, the grades addressed as coarse, are the ones with a lognormal particle size distribution (PSD) with a mean size of 6 µm, while the grades, addressed as fine, are the one with an equivalent PSD with mean size of 0.8 µm.

The effect of the cooling rate on the limits of the C-window was assessed by cooling the samples with various mean grain sizes and C contents, at two different cooling rates, being 15 °C/min and 80 °C/min. From now on denoted as slow (S) and fast (F) cooling rates (Fig. 2).

The second matrix aims to study the effect of the sintering temperature on the formation of η₆ carbides. In this case the cooling rate has been fixed to the slow condition, and two nominal carbon contents representatives of 3-phase and 2-phase microstructures have been studied. These samples vary in grain size, including an additional mean grain size of 0.45 µm and are sintered at 1480 °C and 1440 °C holding temperatures. (Fig. 3)

For the processing of all powder mixtures, 1.5 wt% of paraffin was added and wet-milling was performed in a rotation ball mill with cyclohexane as liquid medium with a weight ratio of milling balls to powder of 4:1. The milling time was 24 hours for the coarse blends and 72 hours for blends with mean size 0.8 and 0.45 µm, respectively. The milled powders were dried in a Rotovapor at 60 °C, and then sieved and pressed at 200 MPa into prismatic green bodies.

Sintering experiments were carried out in a sinter-HIP furnace. All samples were subjected to the same heating cycle, but the sintering temperature and cooling step were modified depending on the desired conditions.

The heating step was characterized by a heating rate of 10 °C/min up to 1250 °C, an isothermal holding for 30 min and a final heating rate of 3 °C/min up to the selected sintering temperature (1480 °C or 1440 °C). The total sintering time for all samples was 30 minutes. For the samples
<figure><img src="image_5.png" /><figcaption>**Fig. 2.** C contents measured for samples sintered at 1480 °C, for fast and slow cooling rates and 6 µm and 0.8 µm sized grades. L and H have the same meaning as in **Fig. 1**. (Targeted compositions are shown as shapes described in the black legend).</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>**Fig. 3.** C contents measured for samples with a target C content of 5.30 and 5.35 wt% C, sintered at 1440 °C or 1480 °C, and cooled at 15 °C/min. L and H have the same meaning as in **Fig. 1**.</figcaption></figure>

cooled at 15 °C/min, a pressure of 6 MPa was built in the sintering chamber and held for the last 10 minutes. For the samples cooled at 80 °C/min, a pressure of 10 MPa was built in the sintering chamber and held for 10 minutes at sintering temperature and along the whole cooling step.

## 2.3. Sample characterization

The resulting carbon content after sintering was measured by means of infrared analysis of the gas products resulting from the combustion of crushed samples in a LECO CS 230 equipment.

All samples were studied through light optical microscopy (LOM) in an as-sintered and Murakami-etched condition. Quantification of phase fractions was performed on etched samples through image analysis at different scales.

Selected samples were analyzed by X-ray diffraction (XRD) in a Bruker D8 Discover, with a LynxEye 1D detector, in Bragg-Brentano geometry using Cu-Kα radiation. The diffraction patterns were recorded from 30 to 47° 2θ with angular steps of 0.02° and a dwell time of 3.5 s per step. Rietveld analyses of XRD spectra, using peaks for WC and η₆-carbide, was done to confirm the quantification results.

Scanning electron microscopy (SEM) was performed on additional selected samples in a Hitachi S3700N FEG-SEM using secondary electron (SE) and back scattered electron (BSE) detectors with acceleration voltage of 15 kV and a current of 7.4 nA. Chemical composition in the binder and in the carbides was semi-quantitatively analyzed by means of point analysis of Energy-Dispersive X-ray Spectroscopy (EDS) patterns. At least 10 points of each of the reported conditions were measured. An acceleration voltage of 15 kV was used for all measurements, and only W, Fe and Ni were considered for the quantification of the composition. L- and M- series were tested for W quantification resulting in comparable results with slightly less scatter for the L-series. L-series were thus used for binder and carbides composition quantification.

The carbides of samples with target C content of 5.15 wt% were studied in detail with a combination of Electron Back-Scattered

Diffraction (EBSD) in SEM and Selected Area Diffraction (SAD) in Transmission Electron Microscopy (TEM). EBSD characterization was performed in a FEI Nova600 FIB-SEM (FEG) with an Oxford Instruments Symmetry CMOS at 20 kV acceleration voltage and 2.4 nA beam current.

TEM samples were prepared as lift-outs in a Zeiss Crossbeam focused ion beam and scanning electron microscope, FIB-SEM. Before ion milling the area was covered by a protective layer of platinum. Ion milling was done at 30 kV for cutting the lamella and attaching it to the Omniprobe tungsten needle. The sample was attached to a Cu-grid by platinum deposition and thinned down to electron transparency with final steps at 100 pA, 10 kV, and 50 pA, 5 kV.

TEM and STEM (scanning transmission electron microscopy) was performed in a FEI Titan Themis 200 with probe corrector and SuperX EDS, operated at 200 kV. SAD was done with a 150 nm in diameter aperture and plane spacing was measured using diffraction patterns of gold-palladium nanoparticles as a reference. The composition was studied by means of transmission EDS using Velox for acquisition and standardless quantification, with Cu and Pt included for deconvolution only.

After the SAD analysis additional EBSD maps were produced with a new phase list for indexing, this time in a Jeol JSM-7800 F FEG-SEM with a Bruker analytical systems e-Flash HR EBSD detector at 15 kV acceleration voltage.

Regarding mechanical properties, Vickers hardness with a load of 30 kg was measured for all samples in a Struers Duramin-40 equipment. The indentation fracture toughness (IFT) was determined from Vickers indentations using 30 kg load for all samples except for coarse samples where 50 kg load was needed to induce suitable crack formations. Palmqvist crack geometry was assumed for the estimation of the IFT in all samples.

# 3. Experimental results

## 3.1. Phase identification

XRD diffraction patterns obtained for samples in matrix I with target C content 5.15 wt% C and 5.25 wt% C are shown in Fig. 4, together with the theoretical diffraction peaks in the analyzed 2θ range [31°, 47°] for the phases expected according to thermodynamic calculations for these compositions and two additional sub-carbides (κ-carbide and W₂C) stable at C contents lower than the nominal one. The spectra are in all cases dominated by the WC diffraction peaks, located at 2θ values of 31.46° and 35.86°. The theoretical peak for the {111} planes for pure Fe-Ni 1:1 FCC alloy is located at 43.6°. The corresponding peak observed experimentally is located at a 2θ value of 42.44±0.034, which indicates an increased lattice parameter as expected due to the dissolution of W in FCC. Also, η₆-phase is detected in the XRD diffraction patterns. The detected peaks correspond well with theoretical spectra for M₆C carbide with composition (W₃M₃C). Regarding the intensity of these peaks, 5.15 wt%C samples show a higher intensity than 5.25 wt%C samples indicating higher volume fractions.
<figure><img src="image_7.png" /><figcaption>**Fig. 4.** XRD patterns obtained for samples with target 5.15 wt%C containing core-rim structures and samples with target 5.25 wt%C only containing dendritic η₆ for processing conditions of matrix 1.</figcaption></figure>

## 3.2. Morphology and composition evolution of precipitates – Effect of C content

LOM micrographs representative of the four types of microstructures observed along the C series for each of the processing conditions in experimental matrix I are presented in Fig. 5. For all cooling rates and initial grain size combinations graphite colonies are present in the microstructures with the highest C contents, showing that the graphite precipitation limit was passed. As the C content decreases, samples between target [5.35–5.425] wt%C in the fine samples and between target [5.35–5.45] wt%C in the coarse samples present a two-phase

microstructure only composed of FCC binder and WC. For both grain sizes, in the C region between target [5.25–5.30] wt%C, η₆-carbides form as dendrites even if not well developed.

For samples with target C content of 5.15 wt% a change in the precipitates morphology is observed, and while some dendrites are still observed, the fraction of more spherical carbides increases substantially. This difference suggests a change in the carbide formation. Further, some differences can be highlighted between the 5.15 wt%C samples and the samples with higher C content containing η₆. Fig. 6 shows SEM micrographs that are representatives of the two types of η₆-carbides obtained regardless of grain size and cooling rate for samples sintered at 1480 °C. Samples with a target composition of 5.15 wt%C show a core-rim structure in which the core is lighter than the rim, and the fraction of engulfed WC grains is lower in the core than in the rim. The rim structure is analogous to the η₆-carbides in samples with target compositions between [5.25–5.30] wt%C, where no core-rim structure is observed and WC grains are trapped by η₆ during its growth.

The volume fraction of WC inside the carbides, as seen in Fig. 7, is determined by measuring the area fraction through image analysis of at least 10 carbides per condition. The fraction of WC grains trapped inside of the carbides is substantially affected by the grain size but not by the cooling rate. For all studied structures, the fraction of engulfed WC is lower for finer grades. The fraction of engulfed WC in the 5.15 wt%C samples rim structures and [5.25–5.30] wt%C η₆-carbides are 35 % and 54 % for fine and coarse grades respectively. These amounts are lower for the core region, reaching values of 25 and 33 % for the fine and coarse grades respectively.

According to the phase diagram, and indicated by the XRD measurements, no other phases than WC, FCC Ni:Fe and η₆ are expected to be present in the materials microstructure. Thus, the appearance of a core-rim structure with a sharp interface in the lowest C samples needs further analysis. Fig. 8 shows qualitatively the composition of the carbides in samples with target composition 5.15 wt%C measured by means of EDS<sub>Kα-Fe</sub> (Fig. 8b) and measured in TEM-EDS (Fig. 8c). Results from both characterization techniques show that the core structure has higher W and C content, but lower Ni and Fe content than the rim. The compositions inside the core and the rim are homogeneous, and no composition profiles have been detected. When comparing the compositions measured in several core-rim structures in samples with different cooling rates and grain sizes no significant differences were found. Since EDS is not adequate to precisely measure light elements, such as C, only the ratio of metallic elements measured in at% and averaged for all processing conditions is reported in Fig. 9. The comparison between core and rim compositions shows a clear trend where the relative content of W is higher in the core, and the lower W content in the rim is at the expense of Fe. Approximately, the amount of W is twice the amount of (Fe+Ni) in the core structure, while in the rim this ratio is approximately 1:1.

In the considered thermodynamic database (TCFE10), η₆ carbides are modeled as the M6C_E93 phase. This phase is described as (Fe, Ni)₂(W)₂(Ni, Fe, W)₂C, and thereby its (W/(Fe+Ni+W)) metallic ratio ranges from 66.7 % to 33.3 %, including all of the measured compositions of the core, the rim and the dendrites. For the case of the core, the measured ratio (66 %) could correspond to η₆ with metastable composition of M₂W₄C. However, such η₆ composition is not predicted by equilibrium at any temperature along the whole processing temperature range. Fig. 9 shows the metallic equilibrium ratios for η₆ calculated for a material with the experimental C content from sintering temperature to T<sub>Freeze-in</sub>. These calculations show that the relative content of W in η₆ increases with temperature, the highest (W/(Fe+Ni)) ratio is reached at the sintering temperature, being 1.17. This metallic ratio is still far from the high values measured in the core. On the other hand, the ratios measured in the rim and the η₆-carbides in samples with C contents between 5.25 and 5.30 wt%C are the same (50 %) and in good agreement with those predicted by equilibrium at the end of the solidification process (ca. 1350 °C).
<figure><img src="image_8.png" /><figcaption>Fig. 5. LOM micrographs representative for the four types of microstructures observed along the processed C series.</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>Fig. 6. BSE-SEM micrographs (x1200) of the carbides observed in the studied samples.</figcaption></figure>

The difference in η₆ morphology observed when comparing samples with target wt%C of 5.15 and [5.25–5.30] in Fig. 5, together with the core-rim structure and its chemical composition variation (Fig. 8) indicate a change in η₆ formation mechanism. The main difference between one microstructure and another may rely on the presence or absence of a third phase other than WC and Liquid when the heating step is completed and the samples reach the sintering holding plateau. For [5.25–5.30] samples that do not present a core-rim structure, η₆ precipitates on the WC-Liq interfaces from the supersaturated liquid as the material is cooled down. Conversely, samples with target C content of 5.15 wt% C may contain stable precipitates already during sintering

(See compositions denoting solubility lines at sintering temperature in Fig. 2).

The more detailed characterization of the core-rim structure by means of EBSD sheds some light on the possible formation of this structure. Fig. 10.a shows the band contrast (BC) from EBSD mapping in a region where both core and rim are measured. It can be seen how the core structure appears as a single coarse grain while the rim structure shows several grains smaller in size. The inverse pole figure (IPF) map in Fig. 10.b further confirms the presence of several grains with distinct orientations in the rim structure. This difference in grain size already hints of a different thermal history for both structures. In the case of the
<figure><img src="image_10.png" /><figcaption>**Fig. 7.** Volume fraction of WC trapped within the carbides for different C contents and mean grain sizes.</figcaption></figure>

core, the time has been long enough and the temperature high enough for coarsening to occur, suggesting that the core regions are already present in the microstructure during the sintering step. The comparison of the IPF measured for coarse 5.15 wt%C samples cooled under fast (Fig. 10.b) and slow (Fig. 10.d) rates also shows the effect of cooling rate on the rim structure. When compared, the rim structures in both samples are polycrystalline, but in the sample cooled under fast conditions,

corresponding to Fig. 10.a and b, more grains in the rim are observed despite the high magnification. Contrarily, in Fig. 10.d, corresponding to the sample cooled under slow conditions, the magnification is lower and still roughly only four grains with distinct orientations are observed in the rim. This shows that faster cooling rates effectively refine the microstructure of the rim.

To further understand the high W contents measured in the core structure, two lamella samples for TEM-analysis were cut across the core-rim boundary and selected area diffraction was performed on both phases. Fig. 11 shows the diffraction patterns obtained for the rim structure from sample 1. The symmetries and plane spacing suggest that this phase is face centered cubic, FCC, η-carbide. The unit cell length as calculated from the diffraction patterns from the first and second lamellas are 11.02 Å and 11.07 Å respectively. Both calculated lattice parameters are in good agreement with previously reported values for the η-phase in Fe based cemented carbides 11.07–11.15 Å [33,34] and also for lattice parameter predicted from first-principle calculation for NiFe₂W₃C, 11.03 Å [35].

Regarding the core, Fig. 12 shows the diffraction patterns obtained from the core structure. These patterns are not a good match with an FCC lattice, showing that the core and the rim are different phases. The six-fold rotation symmetry and suggested large plane spacing of pattern [112]/[1102] together with the orthogonal characteristics and large plane spacing of pattern [100]/[1010] suggest a hexagonal lattice (P6₃/mmc). Tilting to an additional 2 orientations in sample 5 confirms this. The lattice parameters measured from the patterns of sample 1 are a = 7.727 Å and c = 7.783 Å, and the patterns of sample 2 give a = 7.828 Å. The patterns from sample 2 does not have any reflections that can be used for direct determination of c, but calculations based on the value of a gives c = 7.532 Å.

Rautala and Norton describe a hexagonal carbide (MW₃C) in the M-W-C system that they call the κ-phase [36]. Its existence was shown experimentally by Schönberg for the Co-W-C system (a = c =

<figure><img src="image_11.png" /><figcaption>**Fig. 8.** a) BSE-SEM micrograph of fine-grained sample cooled under fast conditions with target composition of 5.15 wt%C showing a core rim structure. b) EDS composition mapping showing the enrichment of Fe towards the rim of the η-carbides. c) Transmission EDS results showing the elemental signal maps for a triple point WC-Core-Rim.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Fig. 9.** Bar charts showing the average metal ratio (%) measured for the core and rim structures in 5.15 wt%C samples and η<sub>6</sub>-carbides in [5.25–5.30] wt%C samples; together with equilibrium metal ratio in η<sub>6</sub>-carbides along temperature for a cemented carbide with 5.24 wt%C.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 10.** a) EBSD-BC and b) EBSP-IPF maps of 5.15 wt%C coarse sample cooled under fast conditions, showing WC grains core and rim. c) EBSD – Phase (overlapping BC) map and d) EBSD-IPF map at lower magnification of 5.15 wt%C coarse sample cooled under slow conditions.</figcaption></figure>

7.826–7.848 Å) [37], by Pollock and Stadelmaier for the Fe-W-C system (a = c = 7.806 Å to 7.810 Å) [34] and by Khomenko et al. for the Ni-W-C system (a=7.834 Å c=7.8048 Å) [38]. Both symmetry and lattice parameters match very well with the core structure. However, the measured lattice parameters in this work are slightly smaller than most of the values reported in literature. This can be explained by the difference in composition observed for the quaternary carbides

characterized in this study in relation to the ternary carbides reported in literature, and the effect that this compositional difference may have on the carbides' crystal structure. The measured composition of the κ-carbides in samples with 5.15 wt%C target composition is Fe₂Ni₂W₈C₃, while the carbide formula proposed in literature for ternary carbides is M₃W₉C₃. The detailed crystallographic characterization performed for F-W-C κ-carbides [39], reveals some of the effects of the composition on
<figure><img src="image_14.png" /><figcaption>Fig. 11. Diffraction patterns from the rim structure.</figcaption></figure>

<figure><img src="image_15.png" /><figcaption>Fig. 12. Diffraction patterns from the core structure.</figcaption></figure>

the crystallographic parameters. In Fe based κ-carbides, despite Fe having a larger atomic radius than Co, these carbides show a smaller lattice parameter than it's Co analogous. This is explained by Fe, in contrast to Co, occupying some of the W lattice sites, and thus reducing the unit cell. In fact, the exchange of one W atom for one Fe atom in the $Fe_3W_9C_3$ carbide barely changes the stability of the carbide [40]. Contrarily, the replacement of one W atom by a Ni atom significantly decreases the carbide's stability. Thus, in the measured carbides, one Fe atom replaces one W atom in the 6 h positions of the $P6_3/mmc$ lattice,

leading to a decrease in unit cell volume. Additionally, Fe and Ni are interchangeable atoms in the 6 h positions, but Ni having smaller atomic radius than Fe. Therefore, Ni occupies two thirds of the Fe-6 h lattice sites, further reducing the unit cell volume and enhancing the observed anisotropy between a and c. TEM characterization has allowed the identification of the core structure as κ-carbide with certainty, thus, the absence of signal from the κ-carbide in the XRD data (see Fig. 4 ) is probably due to the low total amount and the overlap of the diffraction peaks with highest intensity with peaks from WC and $\eta_6$.
Information on the thermal and compositional stability of κ-carbides in ternary systems is limited, and even more so in quaternary systems like the present one. This gap has led κ-carbides not to be included in the thermodynamic databases used in the current work. Still some tentative conclusions can be drawn from the scarce available literature [36,34,33, 41,42]. For all three ternary systems considered (Co, Fe, Ni)-W-C κ-carbide was found to be stable at temperatures below the melting temperature in the approximate compositional space (M,W,C) at%=[(0, 100,0),(0,50,50),(40,46,14)]. The compositional range for ternary κ-carbides is reported to be narrow, however the experimentally measured metallic ratios (see Fig. 9), the fact of κ being a carbide [43] as well as the presence of additional metallic elements in the quaternary systems [42] suggest that the broadening of its homogeneity range is possible. The thermal stability of κ-carbides below the melting temperature is proven in [34,33,41,42] and the possibility of them being stable at temperatures higher than the sintering temperature considered in this work is proposed in [36,43].

The proposed mechanism for the formation of the core-rim structure is illustrated in Figs. 13 and 14, and can be described as follows:

1) During the heating stage, below the peritectic temperature, the material is in a solid powder-like state. At this stage, samples with an understoichiometric composition (with pure W powder added to the blend) contain regions with W contents much higher than the nominal one. In these regions the formation of κ-carbides is locally promoted, and the dispersion of the κ precipitates is directly influenced by the powder dispersion achieved in the blending, milling, and pressing processing steps. At these temperatures, where only solid-state phase transformations may take place and considering data reporting presence of κ at temperatures below the melting temperature, it is reasonable to assume that κ-carbide formed during heating.

2) When the material reaches the melting temperature and liquid forms, the diffusion becomes more efficient, leading to homogenization. Since the nominal W content is not high enough for κ to be stable, κ-carbides will dissolve or transform into η₆ depending on the thermodynamic conditions and thus the composition of the surrounding WC+Liq subsystem. The composition varies with the fraction of metastable κ and in Fig. 13 is shown the carbon content in this WC+Liq subsystem as a function of the fraction of κ. This is obtained through a mass balance considering the measured κ composition and the fraction of WC engulfed in the core (see Fig. 7). The onset of each

of the lines at 0 at% κ is equivalent to the nominal composition since η₆ is not present in the microstructure. The kinetics in the WC+Liq subsystem are considered much faster than those in the WC+κ subsystem, and the transformation of κ is considered to be determined by the equilibrium conditions in the WC+Liq subsystem. Thus, the fraction of metastable κ+WC, will determine if η₆ will be stable or not in the WC+Liq subsystem at the sintering temperature. Fig. 13 shows how higher amounts of κ and lower fractions of trapped WC (fine grades) move the surrounding WC+Liq subsystem towards compositions that are higher in C, inhibiting the formation of η₆. The extent of transformed/dissolved κ at the final stage of sintering is thus conditioned by the amount of initially formed κ and the amount of entrapped WC, and determines the formation of a possible core-rim structure. In samples with target compositions [5.25–5.30] wt %C, the C content of the WC+Liq subsystem is above the phase boundary determining the stability of η₆ for the whole range of κ fractions, except for very small fractions of κ, leading to fully dissolved κ by the time the sintering step is finished. The case for samples with target 5.15 wt%C, presents a different situation and for κ fractions ≤ 3 mol% there is a driving force for the formation of η₆, which could result in a core rim structure. For comparison with experimental results these samples are depicted in Fig. 13 through colored circles, showing the measured κ fraction and sample composition. The arrows approximate the C content of the WC+Liq subsystem of these samples considering that the measured fractions of κ and engulfed WC are representative of those that the system has when reaching the sintering plateau. In case that the κ fraction when reaching the sintering plateau was higher, these samples would still reach the η₆ stability condition as the dissolution of κ progressed. In all samples with target 5.15 wt%C, the WC+Liq subsystem C contents are below or very close to the phase boundary determining the thermodynamic stability of η₆. This condition could lead to κ transforming into η₆ [44], up to an extent ruled by the sintering conditions (time and temperature). This phase transformation is a possible explanation to the core rim structure, where η₆ is observed to cover κ uniformly along its whole perimeter, see Fig. 6. However, this transformation is dependent on diffusion of alloying elements from the liquid and thus its rate will decrease with the thickening of the η₆ layer. The presence of κ-carbides during the sintering holding, promotes the formation of the core-rim structure, in a similar manner as is observed in cemented carbides containing undissolved carbonitrides [45].

<figure><img src="image_16.png" /><figcaption>**Fig. 13.** C-content in the (WC+Liq) subsystem VS κ molar fraction for a) coarse samples, b) fine samples. The colored regions represent the calculated phase stability regions at sintering temperature (1480 °C). Colored dots show the nominal composition and measured κ fraction from samples with target 5.15 wt%C, while the arrows point at the C contents of their respective WC+Liq subsystems.</figcaption></figure>
<figure><img src="image_17.png" /><figcaption>Fig. 14. Schematic showing the different steps involved in the formation of core-rim and dendritic structures.</figcaption></figure>

3) Along the cooling step, the microstructure of the samples in the target composition [5.25–5.35] and 5.15 wt% samples evolve differently. In the former, η₆ carbides nucleate at the WC-Liq interfaces and grow from the melt as predicted by thermodynamics. On the other hand, in the 5.15 wt%C samples, when temperature starts to decrease two factors will promote the nucleation and/or growth of η₆ carbides forming the rim. First, as temperature decreases, the driving force of η₆ increases promoting nucleation and growth. Secondly, the remaining κ will, before covered by a layer of η₆, act as preferential nucleation sites of new η₆ grains, the η₆ grains will continue to grow at the expense of κ, forming the core-rim structure. The proposed evolution of the rim along cooling can be further confirmed since slower cooling rates result in larger η₆ grains and thicker rim compared to fast-cooled samples (see Fig. 10.b and.d).

The understanding of these mechanisms brings valuable information on the transition between dendritic and finely dispersed η-precipitation. It also highlights the importance of further developing the thermodynamic description of cemented carbide systems to include potential intermediate phases, such as κ-carbides or η₁₂, in alternative binder cemented carbide systems. The use of mean field models that consider the material as a system with homogeneous nominal composition present several utilities that will be presented in the following sections. However, the consideration of the powder-solid state inhomogeneities in the model, despite being challenging, would help towards a better understanding on the appearance of intermediate carbides such as κ or W₂C. The formation of such carbides during heating indeed influences the final microstructure. Thus, improved characterization that enable the thermodynamic modeling of these carbides can help turning the appearance of detrimental phases into a useful knob for microstructural design. Understanding that these phases, if properly controlled, can even improve the material properties.

### 3.3. Dynamic C-window

Fig. 15 shows the C-windows for Ni/Fe-1:1-based cemented carbides with 11 wt% binder calculated using the Dynamic C-window model (dashed bars), equilibrium calculation (grey bar) and experimental

<figure><img src="image_18.png" /><figcaption>Fig. 15. Equilibrium (grey bar), experimental (colored bars) and dynamic (Dashed contours) C-windows for the tested conditions. Stars correspond to the experimental compositions shown in Fig. 2. Black and colored stars correspond to samples with three-phase and two-phase microstructure respectively. The red crosses labeled L and H correspond to the limits determined in Fig. 1.</figcaption></figure>

measurements (colored bars). The colored bars are based on the experimental microstructure where colored stars indicate a two-phase microstructure (WC + Binder) and black stars indicate a three-phase microstructure. The third phase in high-C content samples is graphite, while in low-C content samples it is η₆. The equilibrium C-window covers a range between [5.45–5.50] wt% C and is only dependent on the initial composition (red crosses labeled as L and H). While, the experimental and simulated windows can be dependent on the initial WC grain size and cooling rate, besides nominal composition.

The determination of the C-window requires the definition of the limits on both sides of the carbon range. The high-C limit, determined by the precipitation of graphite, is modeled through S-G calculations. The effect of the grain size or the cooling rate on moving the high-C limit towards higher C contents is marginal, so the predicted graphite precipitation limits are close to equilibrium calculations regardless of the
processing conditions. This limited effect is further confirmed by experiments.

On the other hand, when comparing the “Equilibrium” and “Experimental” low-C limits, determined by the precipitation of η, it is clear that the effect of the processing conditions on moving the low-C limit towards lower C- contents (broadening the C-window) is significantly more pronounced than for the graphite precipitation limit. For each of the considered processing conditions, at least one sample show a two-phase microstructure at room temperature even if the three-phase region containing η₆ was crossed during cooling. For coarse series this occurrence increases to three samples, showing the impact of initial grain size on η₆ formation. These occurrences precisely highlight the relevance of defining a dynamic C-window in contrast to using only equilibrium calculations to define the C-window. The presented simulations effectively account for the effect of the initial grain size and the cooling rate on the C-window broadening.

According to the experimental and simulated results (Fig. 15), the tested cooling rates have a limited impact on the broadening of the carbon window, analogously to Co-based cemented carbides [18]. On the other hand, increasing the WC grain size has a considerable effect on shifting the lower C-limit towards lower C contents. For the coarse-grained samples, the C-window is approximately 2.5 times larger than according to equilibrium calculations, increasing from 5.45 to 5.50 wt%C to 5.38–5.50 wt%C. In the case of the samples with a mean size of 0.8 µm, the measured and predicted C-windows are 1.8 times larger than the equilibrium one. The effect of the WC grain size on the C-window broadening observed in Fe-Ni systems, is higher than the one reported for Co-based systems [18]. This may be due to the intrinsic grain growth inhibition effect of Fe maintaining the smaller WC grain size, making the effect of additional nucleation sites of finer grades more relevant.

The ICME nature of the current work allows to use the parametrization obtained from the Fe-Ni 1:1 system to study the larger framework of the Fe-Ni system in general. Fig. 16 presents the equilibrium and the simulated dynamic C-window for cemented carbides in the Fe-Ni system

with 11 and 20 wt% binder content. In Fig. 16 only the two cases of 0.8-Slow and 6-Fast are shown as representatives of the extreme cases for η₆ precipitation conditions. Regarding experimental points from literature, one publication from Uhrenius [46] and one from Prakash [47] are considered. The known processing parameters relevant for the calculation of the Dynamic C-window are listed in Table 2. A cooling rate of 5 °C/min, representative of normal production in-furnace rates, is considered for all calculations of the 20 wt% binder dynamic C-windows. The chosen grain size of 1 µm aligns with Prakash’s data and appears reasonable for Uhrenius’ data. The sintering temperature has been set at 1410 °C for all Fe-Ni ratios below 0.75 and at 1460 °C for higher ratios, to account for the increased melting temperatures for higher Ni contents.

The data reported by Uhrenius [46] is in good agreement with simulations for both Fe-Ni studied ratios. Within these results, the data points reported as “minor η₆” are of special relevance since they help locating the transition region where C content is low enough for η₆ to start precipitating with a higher precision. This may help in assuring with a higher reliability the fact that the Dynamic C-window is nearly double the width of what equilibrium predicts for samples with a Ni:Fe ratio of 0.35. The data reported by Prakash [47], is also in good agreement with simulations, having special relevance samples labeled as

<caption>**Table 2**</caption>

Experimental conditions for considered literature data and conditions considered in calculations shown in Fig. 16.

<table><thead><tr><th></th><th>B. Uhrenius<br>[46]<br>In furnace<br>(Not reported)</th><th>L.J. Prakash<br>[47]<br>In furnace<br>(Not Reported)</th><th>Simulation</th></tr></thead><tbody><tr><td>Cooling Rate [°C/min]</td><td></td><td></td><td>5</td></tr><tr><td>WC grain size [µm]</td><td>Not reported</td><td>[0.9–1.3]</td><td>1</td></tr><tr><td>Sintering Temperature<br>[°C]</td><td>1410</td><td>1400</td><td>1410<br>[0–0.75]<br>1460<br>[0.75–1]</td></tr></tbody></table>

<figure><img src="image_19.png" /><figcaption>**Fig. 16.** Equilibrium and Dynamic C-window for cemented carbides with 20 wt% and 11 wt% Fe-Ni binder system for Ni:Fe ratio from 1 to 0. Samples from current work with Fe-Ni 1:1 (ratio 0.5) have been shifted in Fe-Ni ratio to improve graphic visualization. Data for systems with 20 wt% binder is obtained from ref. [46,47].</figcaption></figure>
AD1.4, AC1.2 and AA1.4 in Fig. 16. For these samples, the consideration of process kinetics was able to explain their two-phase microstructure despite lying outside of the C-window predicted by equilibrium. One exception for the good correlation between simulations and data reported by Prakash can be found in sample AZ0.4. This sample is reported to have a 2-phase microstructure, but precipitation of graphite would be expected from simulation results. The shift of the 2-phase region towards understoichiometric in pure Ni binders is a well-stablished fact, thus the explanation regarding the absence of graphite in sample AZ0.4 remains unclear to the authors of current work.

### 3.4. η<sub>6</sub>-phase size distribution – Effect of sintering temperature and impact on properties

Fig. 17 shows the microstructures obtained for a fixed composition of 5.30 wt% C and a cooling rate of 15 K/min but changing the sintering temperature and the initial grain size. A notable impact on η<sub>6</sub> distribution is observed, where decreasing the sintering temperature from 1480 °C to 1440 °C changes the morphology of η<sub>6</sub> precipitates from large dendrites to finely dispersed particles. For the measured C contents of these samples (see Fig. 3), equilibrium predicts a two-phase material (WC+Liq) at 1480°C sintering temperature, and three-phase material (WC+Liq+η<sub>6</sub>) at 1440°C. This suggests that the final η<sub>6</sub>-carbide dispersion depends on the number of nuclei formed along heating that remain stable at sintering temperatures. For materials with a C content high enough not to form η<sub>6</sub> or κ along the heating step or that are sintered at high enough temperatures to completely dissolve the third phases formed along heating, η<sub>6</sub>-carbides grow dendritically along cooling, reaching big sizes due to the fast kinetics of solidification.

Fig. 18 shows the experimental and calculated number densities for all samples sintered at 1480 °C for different processing conditions. All simulations except those of samples with target 5.15 wt%C consider that no third particles are present at the initial state of the simulation (when the cooling starts). Simulation trends align with the number densities measured experimentally. In both cases, an increase in the cooling rate or a decrease in the initial WC grain size leads to higher number density of η<sub>6</sub>-carbides. However, the effect of increasing cooling rate on the number density is overestimated in the simulations. It can be highlighted that a significant increase in number density, both measured and calculated, is observed when comparing samples with 5.15 wt% C target

composition to [5.25–5.30] wt%C samples. This rise in number density in combination with the core-rim structure observed in 5.15 wt%C samples suggests that the presence of κ-carbides when reaching the sintering plateau helps improving the homogeneous dispersion of the precipitates.

Considering the effect of κ on the dispersion of precipitates in samples sintered at 1480 °C, and the observation of the fine distribution of precipitates in samples sintered at 1440 °C, it is of interest to evaluate if κ is also present in samples sintered at 1440 °C (experimental matrix II). Fig. 19 shows SEM micrographs representative of the η<sub>6</sub>-carbides in these samples. No presence of κ nor the core-rim structure is identified, and a homogeneous average composition of (W<sub>2.7</sub>Fe<sub>2.0</sub>Ni<sub>1.3</sub>)C is measured. This composition lies close to the one for the rim structure and η<sub>6</sub>-carbides measured in high temperature samples, but with slightly lower W and higher Fe and Ni. This is in good agreement with trends of equilibrium composition with decreasing temperature. Despite the absence of core-rim structure, the high number densities in these samples suggest that the nucleation step took place in the solid state [48]. This may have happened through the nucleation of η<sub>6</sub>-carbides during heating and their presence during sintering, or through the formation of small fractions of κ carbides during heating and their complete transformation into η<sub>6</sub> during sintering. Regardless of the transformation route, understanding the effect of the initial WC grain size on the PSD of the precipitates nucleated along heating is important in order to control the structure through processing conditions. Experiments show that decreasing the initial WC grain size leads to a moderate increment of the number density (Fig. 20.a and b) and a decrease in the mean particle size (see Fig. 20.a) of the precipitates. Especially for the coarser samples, the mean size of the η<sub>6</sub> precipitates increases more significantly and a higher irregularity of their shape (round to dendrite transition) starts to be visible (see Fig. 17 and Fig. 19).

Once the effect of the sintering temperature and initial grain size over the η<sub>6</sub> PSD are understood, it is interesting to see how this is translated to the material's mechanical properties. Fig. 21.a and Fig. 21.b show the hardness measured for the entire C-series in experimental matrix I for coarse and fine samples respectively. Both sizes show a comparable trend in hardness where increasing volume fraction of η<sub>6</sub>-carbides results in higher hardness; the hardness of two-phase samples shows a plateau; and the increasing volume fraction of graphite results in lower hardness. The observed trends for hardness as a function of the

<figure><img src="image_20.png" /><figcaption>Fig. 17. Microstructures obtained for samples with target composition of 5.30 wt% C, for initial WC grain sizes of 0.45, 0.8 and 6 µm, cooled at 15 K/min and sintered at 1480 and 1440 °C.</figcaption></figure>
<figure><img src="image_21.png" /><figcaption>Fig. 18. Experimental and simulated precipitate number densities for samples sintered at 1480°C.</figcaption></figure>

<figure><img src="image_22.png" /><figcaption>Fig. 19. η<sub>6</sub>-carbides in a) 6 µm b) 0.8 µm samples with target composition 5.40 wt%C sintered at 1440 °C.</figcaption></figure>

<figure><img src="image_23.png" /><figcaption>**Fig. 20.** a) PSD of η<sub>6</sub>-carbides in 5.30 wt%C samples sintered at 1440 °C b) Number density of η<sub>6</sub>-carbides in 5.30 wt%C samples sintered at 1440 °C and 1480 °C for different initial grain sizes.</figcaption></figure>

C content (thus function of the microstructure) are similar for all cooling
rates and initial grain sizes. However, it can be highlighted that the
hardness increase in the fine samples resulting from having η<sub>6</sub> present in
the microstructure is significantly sharper than in the coarse samples.
This can be due to the finer dispersion (higher number densities)

obtained for the η₆ precipitates in the fine samples.

While the increase in hardness due to η₆ presence has been consis-
tently reported in literature [46,48,49,50], it often compromises the
material toughness. In this sense it is worth studying if such compromise
can be damped by refining the η₆ precipitates. Fig. 22 shows the results
<figure><img src="image_24.png" /><figcaption>Fig. 21. Hardness measured for samples sintered at 1480°C belonging to experimental matrix I.</figcaption></figure>

obtained from Vickers indentation testing of all samples in experimental matrix II. From Fig. 22.a it is observed that all samples, except those with big dendritic η₆-carbides, show properties within the state of the art (SOTA) hardness-toughness combination. For the case of samples with coarse η₆ dendrites, their hardness falls within the range expected for cemented carbides with their grain size and binder content, but their indentation fracture toughness (IFT) is lower than SOTA cemented carbides. A more straight forward interpretation of such results can be seen in Fig. 22.b and c, where the hardness for samples with 3 and 2-phase microstructures are very close to each other, and the hardness of samples with a finely dispersed 3-phase microstructure show a slight increase. When studying the effect of microstructure on the IFT, 2-phase samples are those reporting the highest fracture toughness with no significant effect of the sintering temperature. The presence of η₆-carbides in the microstructure decreases the fracture toughness for both sintering temperatures. This can be explained by the growth of η₆ (a

brittle carbide) at the expense of the binder (a tough metal). However, Fig. 22.c shows that despite having similar η₆-volume fractions (see Fig. 17), the decrease in fracture toughness is lower in samples sintered at 1440 °C with finely distributed precipitates. For these samples, the hardness is the highest among all the samples in matrix II, and the fracture toughness lies in between the 2-phase samples and the 3-phase coarse-η₆ samples.

Some additional caution needs to be considered when discussing the IFT results. For indentations made in 2-phase samples and finely distributed 3-phase samples (T<sub>Sint</sub>=1440 °C), the measured IFT is independent of the location of the indentation and the cracks shape and arrangement are suitable to consider Palmqvist crack geometry. Conversely, the IFT measured in 3-phase samples sintered at 1480 °C is strongly affected by the indentation location, and cracks arrangement diverges from the ideal Palmqvist assumption. In these samples, IFT values equal to those of two-phase materials can be measured if the

<figure><img src="image_25.png" /><figcaption>Fig. 22. Property chart of state-of-the-art materials [52] with overlapped hardness and toughness values measured for samples in matrix II. b) Measured hardness and c) measured toughness.</figcaption></figure>
indentations are located far from the big η₆-dendrites. Otherwise, when the indentations are located close to the coarse precipitates, the cracks are unevenly distributed and progress up to significantly longer distances in those indentation corners located close to a η₆ dendrite. The IFT reported from these measurements considers the addition of the four crack lengths formed on the indentation corners as recommended by Roebuck [51]. The inhomogeneous nature of these microstructures makes the determination of bulk Kᵢₚ from IFT challenging. If the crack length developed in the corners containing η₆ are considered to be the one representative for the determination of Kᵢₚ (weakest-link approach), the obtained IFT values would be lower than the ones reported in the current work. On the other hand, if one only considered cracks formed in the two-phase region, the measured Kᵢₚ values would be in the regime measured for two-phase samples. Thus, the reported values should be taken with caution and interpret them qualitatively rather than quantitatively. If more precise Kᵢₚ values for such samples was needed, experimental techniques involving bigger testing scales such as sample bending, or pre-notched bar testing methods would be required.

The presented results opens up the door to consider η₆-carbides as third phases which microstructure can be tweaked to reach a certain combination of desired properties. In this regard, the precipitation of η₆ carbides is controlled rather than just avoided as in the traditional cemented carbides processing. The formation of finely distributed η₆ precipitates along the heating stage, could be retained by exploring C contents slightly lower than the Dynamic C-window low-C limit and keeping the sintering temperature just above the peritectic temperature to avoid the growth of solid-state formed η₆ during sintering.

## 4. Conclusion

In conclusion, an extensive study on the processability of Fe-Ni 1:1 based cemented carbides with special focus on the determination of the C-window considering kinetic effects is presented. Based on an ICME framework, the presented analysis combines experimental and modeling work to reach conclusions on systems with Fe-Ni 1:1 11 wt% binder and extrapolate them to other binder contents and Fe-Ni ratios. It can be concluded that for Fe-Ni based cemented carbides:

* The width of the C-window is effectively widened, compared to equilibrium calculations, by considering the kinetics involved in the formation of the detrimental phases limiting the C-window.

* Increasing the cooling rate from 15 K/min to 80 K/min resulted in very mild broadening of the measured C-window, ≈6% broader.

* Increasing the initial WC grain size from 0.8 to 6 µm resulted in a more significant C-window broadening, being 33% broader for coarse grades and 110% broader when compared to equilibrium calculations.

A detailed examination has been dedicated to study the formation mechanism of η₆-carbides, reporting, for the first time to our knowledge, the significance of core-rim structures formation at substantially understoichiometric compositions. The formation mechanism proposed for such structure, involves the formation of κ-carbides along the heating step. These carbides are currently not included in the available thermodynamic databases, but according to several experimental studies [36; 34; 33; 41; 42] at the considered sintering temperatures and at nominal compositions around the C-window they are expected to be metastable. The lack of thermodynamic description of these carbides in the used databases makes their detailed modeling complex. However the observed influence of their presence on the morphology and distribution of η₆ carbides, suggests that further studies may help to tailor the material's properties when designing new cemented carbide grades. For instance, the importance of precipitating η₆ carbides from the melt or having third phases present when reaching the sintering step is discussed. In this sense, experiments at lower sintering temperatures where η₆-carbides are present during the sintering step showed significant

changes in the obtained microstructures of 3-phase materials that could be the subject for further investigation of cemented carbides properties.

## Funding sources

This research did not receive any specific grant from funding agencies in the public, commercial, or not-for-profit sectors.

## CRediT authorship contribution statement

**Victor Lamelas Cubero:** Writing – review & editing, Writing – original draft, Visualization, Software, Methodology, Investigation, Formal analysis, Conceptualization. **Lisa Toller-Nordström:** Writing – review & editing, Investigation. **Manon Bonvalet Rolland:** Writing – review & editing, Supervision, Software, Conceptualization. **Martin Walbrühl:** Writing – review & editing, Supervision. **Raquel de Oro Calderón:** Resources, Investigation. **Annika Borgenstam:** Writing – review & editing, Supervision, Funding acquisition.

## Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Data Availability

Data will be made available on request.

## References

[1] C.M. Fernandes, A.M.R. Senos, Cemented carbide phase diagrams: a review, Int. J. Refract. Met. Hard Mater. Vol.29 (2011) 405–418, https://doi.org/10.1016/j.ijrmhm.2011.02.004.

[2] Roulon, Zoë, Effect of binder on sintering and microstructure of cemented carbides, S. I. Univ. Grenoble Alpes (2019). NNT: 2019GREAI061.

[3] M. de Nicolás, et al., Design of alternative binders for hard materials, Int. J. Refract. Met. Hard Mater. Vol. 87 (2020) 105089, https://doi.org/10.1016/j.ijrmhm.2019.105089.

[4] S.-H. Chang, S.-L. Chen, Characterization and properties of sintered WC–Co and WC–Ni–Fe hard metal alloys, J. Alloy. Compd. Vol 585 (2014) 407–413, https://doi.org/10.1016/j.jallcom.2013.09.188.

[5] C.M. Fernandes, A.M.R. Senos, M.T. Vieira, Control of eta carbide formation in tungsten carbide powders sputter-coated with (Fe/Ni/Cr), Int. J. Refract. Met. Hard Mater. Vol. 25 (2007) 310–317, https://doi.org/10.1016/j.ijrmhm.2006.07.004.

[6] R. de Oro Calderon, C. Edtmaier, W.-D. Schubert, Novel binders for WC-based cemented carbides with high Cr contents, Int. J. Refract. Met. Hard Mater. Vol.85 (2019) 105063, https://doi.org/10.1016/j.ijrmhm.2019.105063.

[7] European Chemical Agency. An agency of the European Union. echa.europa.eu. [Online] [Cited: 06 05, 2023.] https://echa.europa.eu/substance-information/-/substanceinfo/100.028.325).

[8] European Comission. Critical Raw Materials Resilience: Charting a Path towards greater Security and Sustainability. European Comission Website. [Online] 09 02, 2020. [Cited: 05 26, 2023.] https://ec.europa.eu/docsroom/documents/42849). COM(2020) 474.

[9] M. Walbrühl, et al., A new hardness model for materials design in cemented carbides, Int. J. Refract. Met. Hard Mater. Vol 75 (2018) 94–100, https://doi.org/10.1016/j.ijrmhm.2018.04.004.

[10] R. Schaller, et al., Mechanical behaviour of WC–11 wt% Co studied by bending tests, internal friction and electron microscopy. J. Hard Mater. Vol. 3 (1992) 351–362.

[11] Strelsky, J. García, W, Chapter 7: Process development and scale up of cemented carbide production. [book auth.] Maximilian Lackner. Scale-up in Metallurgy. s.l, Verlag Process Eng Engineering GmbH, 2010.

[12] A.F. Guillermet, An assessment of the Fe-Ni-W-C phase diagram, Z. Met. Vol.78 (1987) 3.

[13] R. de Oro Calderon, et al., Fundamental aspects of phase formation in WC-based cemented carbides containing FeMn-based binders, Int. J. Refract. Met. Hard Mater. Vol. 117 (2023) 106411, https://doi.org/10.1016/j.ijrmhm.2023.106411.

[14] P. Zhou, et al., Experimental investigation and thermodynamic assessment of the C–Co–Fe–Ni–W system, Int. J. Refract. Met. Hard Mater. Vol. 54 (2016) 60–69, https://doi.org/10.1016/j.ijrmhm.2015.07.021.

[15] R. Steinlechner, A study on WC-Ni cemented carbides: constitution, alloy compositions and properties, including corrosion behaviour, Int. J. Refract. Met. hard Mater. Vol.103 (2022) 105750, https://doi.org/10.1016/j.ijrmhm.2021.105750.
[16] W.D. Schubert, et al., Aspects of sintering of cemented carbides with Fe-based binders, Int. J. Refract. Met. hard Mater. Vol. 49 (2015) 110–123, https://doi.org/10.1016/j.ijrmhm.2014.07.028.

[17] L. Kaufman, H. Bernstein, Computer calculation of Phase Diagrams, Academic Press, New York, 1970. ISBN: 012402050X.

[18] V. Lamelas, et al., Modelling the formation of detrimental phases in cemented carbides, Mater. Des. Vol. 228 (2023) 111823, https://doi.org/10.1016/j.matdes.2023.111823.

[19] Y. Gao, et al., Effect of Fe/Ni ratio on the microstructure and properties of WC-Fe-Ni-Co cemented carbides, Ceram. Int. Vol. 44 (2) (2018) 2030–2041, https://doi.org/10.1016/j.ceramint.2017.10.148.

[20] B. Wittmann, W.-D. Schubert, B. Lux, WC grain growth and grain growth inhibition in nickel and iron binder hardmetals, Int. J. Refract. Met. Hard Mater. Vol. 20 (2002) 51–60, https://doi.org/10.1016/S0263-4368(01)00070-1.

[21] Y. Wang, et al., Microstructure evolution in the cemented carbides WC–Co I. effect of the C/W ratio on the morphology and defects of the WC grains, Phys. Status Solidi (a) Vol. 193 (2) (2002) 271–283, https://doi.org/10.1016/j.matdes.2023.111823.

[22] I. Borgh, et al., Effect of carbon activity and powder particle size on WC grain coarsening during sintering of cemented carbides, Int. J. Refract. Met. Hard Mater. Vol. 42 (2014) 30–35, https://doi.org/10.1016/j.ijrmhm.2013.10.004.

[23] N.A. Ozbek, et al., Application of deep cryogenic treatment to uncoated tungsten carbide inserts in the turning of AISI 304 stainless steel, Met. Mater. Trans. A Vol. 47A (2016) 6270–6280. DOI: 10.1007/s11661-016-3767-1.

[24] Vinay Varghese, M.R. Ramesh, D. Chakradhar, Influence of deep cryogenic treatment on performance of cemented carbide (WC-Co) inserts during dry end milling of maraging steel, J. Manuf. Process. Vol. 37 (2019) 242–250, https://doi.org/10.1016/j.jmapro.2018.11.030.

[25] C.-H. Xie, et al., Effects of deep cryogenic treatment on microstructure and properties of WC–11Co cemented carbides with various carbon contents, Trans. Nonferrous Met. Soc. China Vol. 25 (9) (2015) 3023–3028, https://doi.org/10.1016/S1003-6326(15)63929-2.

[26] M. Padmakumar, D. Dinakaran, A review on cryogenic treatment of tungsten carbide (WC-Co) tool material, Mater. Manuf. Process. Vol. 36 (2021) 637–659, https://doi.org/10.1080/10426914.2020.1843668.

[27] J.O. Andersson, et al., Thermo-Calc and DICTRA, computational tools for materials science, Calphad Vol. 26 (2002) 273–312, https://doi.org/10.1016/S0364-5916(02)00037-8.

[28] Thermo-Calc Software. TCFE Steels/Fe-alloys Database version 10. (https://thermocalc.com/products/databases/steel-and-fe-alloys/).

[29] Thermo-Calc Software. MOBFE Steels/Fe-Alloys Mobility Database version 5. (http://thermocalc.com/products/databases/steel-and-fe-alloys/).

[30] M. Bonvalet, et al., Modeling of precipitation kinetics in multicomponent systems: application to model superalloys, Acta Mater. Vol. 100 (Vol. 100) (2015) 169–177.

[31] M. Bonvalet Rolland, A. Borgenstam, Modeling precipitation kinetics in multicomponent alloys during deformation, Front. Mater. Vol. 9 (2022) 958472, https://doi.org/10.3389/fmats.2022.958472.

[32] E. Scheil, Bemerkungen zur Schichtkristallbildung, Z. F. üR. Metallkd. Vol. 34 (1942) 70–72.

[33] M. Bergstöm, The eta-carbides in the ternary system Fe-W-C at 1250 °C, C. Mater. Sci. Eng. Vol. 27 (1977) 257–269.

[34] C.B. Pollock, H.H. Stadelmaier, The eta carbides in the Fe–W–C and Co–W–C systems. Met. Trans. Vol. 1 (1970) 767–770.

[35] D.V. Suetin, N.I. Medvedeva, First-principles study of structure, magnetic properties, and stability of η-Carbides (M,Fe)3W3C (M ¼ Ti, V, Cr, Mn,Co, and Ni), Phys. Status Solidi B Vol. 256 (2019) 1900108. DOI: 10.1002/pssb.201900108.

[36] P. Rautala, J.T. Norton, Tungsten–cobalt–carbon system. Trans. AIME Vol. 194 (1952) 1045–1050.

[37] N. Schönberg, The structure of the Co3W9C4, Acta Metall. Vol. 2 (1954).

[38] N. Yu, Khomenko, et al., X-ray diffraction study of a new phase in the Ni-W-C system, Inorg. Mater. Vol. 56 (6) (2020) 572–576. DOI: 10.1134/S0020268520060072.

[39] A. Hårsta, S. Rundqvist, J.O. Thomas, A neutron powder diffraction study of K-phase in the Fe-W-C system, J. Solid State Chem. Vol. 49 (1983) 118–122, https://doi.org/10.1016/0022-4596(83)90223-2.

[40] I.I. Chuev, et al., Density functional theory calculations of the stability and statistical disorder in crystals of the kappa phase of Me3+x W10-x C3+y (Me = Fe, Co, Ni), Russ. J. Phys. Chem. A Vol. 94 (7) (2020) 1369–1374. DOI: 10.1134/S0036024420070080.

[41] M.-L. Fiedler, H.H. Stadelmaier, The ternary system nickel-tungsten-carbon, Int. J. Mater. Res. Vol. 66 (1975) 402, https://doi.org/10.1515/ijmr-1975-660704.

[42] H.H. Stadelmaier, C. Suchjakul, Overview of the quaternary system iron-nickel-tungsten-carbon, Int. J. Mater. Res. Vol. 76 (3) (1985) 157, https://doi.org/10.1515/ijmr-1985-760301.

[43] A. Hårsta, S. Rundqvist, The crystal chemistry of Kappa-Phases, J. Solid State Chem. Vol. 70 (1987) 210–218.

[44] X.-C. Deng, K.-F. Wang, G.-H. Zhang, Effects of oxide addition on structure and properties of WC-10Co cemented carbide obtained by in situ synthesized powder, Int. J. Appl. Ceram. Technol. Vol. 19 (4) (2022) 1916–1927. DOI: 10.1111/ijac.14058.

[45] H.-O. Andrén, Microstructure development during sintering and heat-treatment if cemented carbides and cermets, Mater. Chem. Phys. Vol. 67 (2001) 209–213, https://doi.org/10.1016/S0254-0584(00)00441-7.

[46] B. Uhrenius, H. Pastor, E. Pauty, On the composition of Fe-Ni-Co-WC-Based cemented carbides, Int. J. Refract. Met. Hard Mater. Vol. 15 (1997) 139–149, https://doi.org/10.1016/S0263-4368(96)00023-6.

[47] Prakash, L.J. Weiterentwicklung von Wolframcarbid Hartmetallen unter Verwendung von Eisen- Basis-Bindelegierungen. Karlsruhe: Kernforschungszentrum Karlsruhe, 1980. KfK 2984.

[48] P. Zhou, Y. Du, W. Lengauer, Morphology of η phase in cemented carbides with Fe-based binders influenced by carbon content and nitrogen atmosphere, Ceram. Int. Vol. 45 (2019) 20774–20779, https://doi.org/10.1016/j.ceramint.2019.07.063.

[49] Y. Tang, et al., Effect of a slight change in carbon content near the upper/lower limit on the microstructure and mechanical properties of WC-10Co cemented carbides, Int. J. Refract. Met. Hard Mater. Vol. 96 (2021) 105465, https://doi.org/10.1016/j.ijrmhm.2020.105465.

[50] C. Qian, et al., Effect of the carbon content on the morphology evolution of the η phase in cemented carbides with the CoNiFeCr high entropy alloy binder, Int. J. Refract. Met. Hard Mater. Vol. 102 (2022) 105731, https://doi.org/10.1016/j.ijrmhm.2021.105731.

[51] B. Roebuck, W. Coles, The Measurement of Palmqvist Toughness for Hard and Brittle Materials, in: Measurement Good Practice Guide, 9, National Physical Laboratory, Teddington, UK, 2008. ISSN 1368–6550.

[52] A. Chychko, et al., HV-KIC property charts of cemented carbides: a comprehensive data collection, Int. J. Refract. Met. Hard Mater. Vol. 103 (2022) 105763, https://doi.org/10.1016/j.ijrmhm.2021.105763.