Spark plasma sintering influence on microstructure and mechanical properties of Ti:Ta/carbonitride ceramic matrix composites

Spark plasma sintering influence on microstructure and mechanical properties of Ti:Ta/carbonitride ceramic matrix composites Jos´ e Manuel C´ ordoba Gallego Departamento de Química Inorg´ anica, Facultad de Química, Universidad de Sevilla (US), C/Tramontana s/n, 41012 Sevilla, Spain ARTICLE INFO Keywords: Ceramic matrix composite Spark plasma sintering Microstructure Mechanical properties ABSTRACT A mechanically induced self-sustaining reaction was carried out to synthesize a Ti 0.9 Ta 0.1 C 0.5 N 0.5 /Co powdered cermets, and then they were sintered by spark plasma sintering. Microstructural parameters effects studied by image analysis, and chemical composition (studied by Rietveld analysis) on the microhardness, hardening rate, fracture toughness, transverse rupture strength, and Young’s modulus were related to the sintering conditions. The optimization of the sintering conditions (1150 ◦C, 30 MPa, and 8 min’ dwell time) drove to a homogeneous microstructure and outstanding mechanical properties. Also, the tantalum was suggested to influence the interfacial energies of the system, yielding a stronger hard phase skeleton. 1. Introduction Technically and commercially, cermets are among the most important composite materials due to their exceptional hardness and wear resistance. Because the constituent phases of titanium carbonitride are exceptionally hard and stable, the cermets based on this material have extraordinary resistance to wear and creep. They are excellent options for use as cutting tools because of these qualities. Carbides like TaC, NbC, Mo 2 C, and WC, are frequently utilized in milling applications and interrupted cuts to enhance cutting performance. These carbides form solid-solution phases of mixed carbonitrides and exhibit remarkable hot hardness and thermal shock endurance. [1–3]. The hard component of cermets, called the master alloy, has a composition that determines both the production conditions and the intended application’s technical performance, making it crucial [4]. It is widely accepted in literature [5–7] that the formation of carbonitride solid solutions in terms of ternary or quaternary component structure is given excellent performance to cermets synthesized from the mixture of unalloyed hard components with Co or Ni as binder. Carbonitride solid solutions, which combine the carbide and nitride components into a single phase, have been recommended as a way to enhance the characteristics of cermets [8]. Based on Ti y M T1-y C x N 1-x carbonitrides (M T = Ta, Nb, Zr, V) with a narrow size distribution and submicrometric character (250 nm), mechanically induced self-sustaining reaction (MRS) has recently been proposed as a trustworthy and straightforward method of producing high purity complex solid solution powders [9–12]. Conventional ceramic material sintering methods involve high temperatures. A high degree of sintering is achieved by heating the green specimen to high temperatures for prolonged periods in a typical sintering technique. The sintered item usually contains large grains as a result of these unfavorable conditions, which regrettably encourage unwanted carbide grain development in the presence of metal liquids. As a result, the tool’s mechanical capabilities are reduced. It is commonly known that a finer microstructure results in better mechanical qualities and a longer product lifespan. Spark plasma sintering (SPS), in comparison to pressureless sintering, is a novel sintering technique that applies pressure and electric current to the sample simultaneously [13–14]. At comparatively low temperatures (aprox. 500 ◦C lower than pressureless sintering), and faster heating rates (aprox. 100 ◦C/min), it enables the rapid consolidation of powder compacts to a high density. The SPS technique has shown promising results in reducing the sintering temperature while simultaneously controlling grain growth and secondary phase formation [15–17]. It is unique for researching bulk nanomaterials because of the quick heating rate and potential to maintain the nanoscale during the rapid sintering cycle. This study looked into how the SPS settings affected the sintering of powders with submicrometric grain sizes, Ti 0,9 Ta 0,1 C 0,5 N 0,5 /Co, that were produced via MSR. The resulting cermets’ microstructure and mechanical characteristics (hardness, fracture toughness, transverse rupture strength (TRS), and Young’s modulus) were assessed and contrasted with those of the comparable cemented carbides and ordinary cermets. E-mail address: [email protected]. Contents lists available at ScienceDirect International Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM https://doi.org/10.1016/j.ijrmhm.2025.107094 Received 4 January 2025; Received in revised form 29 January 2025; Accepted 12 February 2025 International Journal of Refractory Metals and Hard Materials 128 (2025) 107094 Available online 13 February 2025 0263-4368/© 2025 The Author. Published by Elsevier Ltd. This is an open access article under the CC BY license ( http://creativecommons.org/licenses/by/4.0/ ). 2. Materials and methods The Ti 0.9 Ta 0.1 C 0.5 N 0.5 solid solution was synthesized by MRS, and the description of the synthesis procedure and characterization (XRD, SEM, TEM, and image analysis) of the as-synthesized powders performed in a previous research [12]. The obtained powders were sintered using SPS-1050-CE (Dr. Sinter, Japan) at 1150 ◦C under different holding times and pressures (see Table 1). X-ray diffraction diagrams and Scanning Electron Microscopy images of cermets polished surfaces were obtained using similar parameters as described in [18] The lattice parameters, phase volume fraction, and ceramic stoichiometry were determined from the XRD diagram using the Rietveld analysis by the use of the Fullprof [19] computer program. The Vickers test was carried out in a Microhardness FM-700 (FutureTech. Corp) with a load of 49 mN, 196 mN, and 981 mN and dwell time ranging from 5 s to 90 s at room temperature. Several microindentations (12 times) at different locations were taken on the polished cermet for every run. The stated value was found from the average of the measured values ( σ = ±5 %). The fracture toughness (KIC) was measured by using three indentation methods under an indentation load of 1 kgf and they were described below. The pieces for flexural tests were cut from the spark plasma sintered cylinders and built into ingots with bar shapes (dimensions of 3×4×25 mm). The room temperature (RT) flexural strength was carried out in a three-point bending arrangement, using a crosshead speed of 0.5 mm/ min and span of 20 mm (AGS-X 10kN Series Dual Column Electromechanical Test Frame, Shimadzu, Japan). All samples were ground to dimensions using resin-bound diamond grinding wheels and surface grinders to achieve dimensional accuracy as well as good surface finishing. The image analysis was carried out by using the Image Pro-Plus 6.0 software and microstructural parameters were obtained as described in C´ ordoba et al. [20]. 3. Results and discussion 3.1. Compositional analysis The phase composition of the sintered bodies was investigated by XRD analysis as shown in Fig. 1. The XRD patterns show the peaks characteristic for the ceramic phase (•) and the metallic binder phases consisting of the intermetallics (*) Ti x Ta 1-x Co 2 [Fd3m] and (○) Ti x Ta 1x Co [Pm3m] produced during the high-temperature sintering. Table 2 describes the quantification of the phases observed in the XRD patterns, the lattice parameters, the coherent diffraction domains, and the corresponding ceramic composition obtained from Rietveld analysis. As shown in Table 2, the Ti/Ta ratio was somewhat constant regardless of the sintering conditions used. However, the C/N ratio of (Ti, Ta)(C, N) decreased when the sintering conditions were stronger (C/ N ratio; SPS1, 1.50, SPS2, 1.38; SPS3, 0.92; SPS4, 0.75). This carbon loss was attributed to the carbon dissolution increasing in the binder phase. The most significant change in the binder phase content (Table 2) is observed in sample SPS4, when a pressure of 60 MPa was used. Because the powder container of the SPS is not sealed, there was some loss of the binder during the process, and a smaller amount was observed and quantified by the compositional analysis in comparison with the other samples. This fact was corroborated by the presence of solidified droplets at the outside of the container. In Fig. 2 are shown characteristic sample SEM images sintered under Table 1 Summary of spark plasma sintering conditions. Sample Nominal Composition Weight [g] Sintering Parameters 1 Ti 0.90 Ta 0.10 C 0.5 N 0.5 +20 %Co 4 1150 ◦C 30 MPa 120 s 2 12 1150 ◦C 30 MPa 120 s 3 12 1150 ◦C 30 MPa 480 s 4 12 1150 ◦C 60 MPa 120 s Fig. 1. XRD patterns of the sintered cermets by SPS. (•) Ti x Ta 1-x C y N 1-y [Fm3m], (*) Ti x Ta 1-x Co 2 [Fd3m], (○) Ti x Ta 1-x Co [Pm3m]. Table 2 Composition, lattice parameter, coherent diffraction domain of the ceramic phase, and description of every binder phase found by Rietveld analysis. For every phase the % vol. is included. The last column illustrates the goodness of fitness for every analysis carried out. Ceramic Phase Binder Phase Sample Ceramic Composition a (Å) D [nm] % vol. Ceram Co 2 Ti(Ta) [Fd3m] [%] CoTi(Ta) [Pm3m] [%] χ 2 SPS1 Ti 0.90 Ta 0.10 C 0.60 N 0.40 43,162 33 78,7 15,7 5,6 1,8 SPS2 Ti 0.86 Ta 0.14 C 0.58 N 0.42 43,184 67 75,2 19,0 5,8 2,2 SPS3 Ti 0.91 Ta 0.09 C 0.48 N 0.52 43,157 74 77,1 18,0 4,8 3,0 SPS4 Ti 0.92 Ta 0.08 C 0.43 N 0.57 43,170 52 86,5 11,0 2,5 2,7 Fig. 2. SEM images showing the characteristic microstructure of the sintered powders. J.M. C´ ordoba Gallego International Journal of Refractory Metals and Hard Materials 128 (2025) 107094 2 different conditions. Data obtained from the image analysis are presented in Table 3. Quantification of the phases presented on the sintered samples through image analysis gave results consistent with those obtained by the XRD results by Rietveld analysis. All the samples presented a highly dense (densification of ~98 % ±2 %, from image analysis, five images at similar scale were analyzed) and uniform microstructure. Although early milling products contained typical particle aggregates, no anomalous development of grains was detected. This indicates that the sintering between aggregates and fine particles was prevented by the spark plasma sintering. Most of the micrometric particles observed in samples SPS1 and SPS4 displayed a typical core-rim microstructure providing evidence of grain growth by dissolution-precipitation processes despite the low temperature and the short residence time at the maximum sintering temperature. In contrast, the core-rim microstructure appreciated in samples SPS2 and SPS3 is barely visible. The sample SPS1 showed the highest average ceramic particle size (Ø =2.56 μ m), while the samples SPS2, SPS3, and SPS4 were lower, and in crescendo from SPS2 to SPS4 (Ø =1.10, 1.95 and 2.28 μ m, respectively). It is common to use an optical pyrometer focused on the outer surface of the die to measure and control the temperature during SPS experiments. This isn’t good enough for high heating rates, where the SPS process’s viability is one of its most evident benefits. It might be sufficient for very slow heating rates where thermal equilibrium can be reached. Moreover, because heating is resistive, both the sample’s and the die’s conductivities affect the temperatures inside and outside the sample. There is a noticeable temperature difference between the samples when a low-conductive material is utilized (in this example, a composite made of ceramic and intermetallic phases) [21]. Due to the increment in mass on the samples from SPS1 to the rest, the distance to be traveled by the current is increased, and therefore the resistivity of the sample (resistivity quantifies how strongly the material opposes the flow of electric current and it is a function of the distance roamed). So, the jump from 4 g (SPS1) to 12 g (SPS2, 3, and 4) increased the temperature gradient, therefore the temperature reached inside sample SPS1 is higher than in the other ones, and this effect is shown in the coarsening of the SPS1 ceramic grains respect the other samples. Also, the formation of a clear core-rim microstructure in sample SPS1 (Fig. 2 (a)) corroborated the inside higher temperature reached. To form a corerim microstructure, a liquid phase sintering has to be reached during the process, and therefore higher sintering temperatures are needed. When a similar weight was used (samples SPS2, 3, and 4) and similar temperature inside the sample can be supposed, the grain coarsening was due to an enhancement of the element diffusion during the sintering because of the increment in the sintering time (SPS3) and pressure (SPS4). Similar behavior was observed in terms of contiguity and mean free path. Sample SPS1 showed the highest contiguity and lowest mean free path values of the set of samples studied, while samples SPS2, 3, and 4 showed an increment in the contiguity from SPS2(0.25) <SPS3(0.28) < SPS4(0.32) and a decrease in the mean free path from SPS2(1.32) > SPS3(1.27) >SPS4(1.20). As above, the coarsening of the grains due to the higher temperature in sample SPS1 increased the contact between the ceramic particles and therefore the increase in the contiguity and the reduction of the mean free path. 3.2. Mechanical properties In Table 3 is showed the hardness, the fracture toughness, the transverse rupture strength and the young’s modulus of the samples studied. Fig. 3 depicts SEM micrographs showing indentations created on the surface of the investigated object, as well as cracks emerging from them, under 1 kgf force. Hardness values were computed using the dimensions of the impression created on the surface of the specimens, and the values reported in Table 3 are the average of twelve runs. The SPS4 sample showed the highest hardness (16.0 GPa) while the rest of the samples presented a microhardness of about ~14 GPa. This can be explained by the loss of binder during the sintering procedure which made the proper combination of ceramic particle size and contiguity to reach the highest microhardness in sample SPS4. Microindentations at different loads (49 mN, 196 mN, and 981 mN) were performed at several time runs to establish the strengthening of the material by work hardening, and results are presented in Fig. 4. As a result of dislocation movements and dislocation generation within the material’s crystal structure, it is evident that the microhardness increases with run time. This resistance to dislocation-formation shows up as a resistance to plastic deformation, which accounts for the observed Table 3 Vickers hardness, contiguity and average free pathway, with the average ceramic particle size, ceramic volume, binder volume and porous volume that were extracted from the image analysis of the selected samples. Image Analysis Mechanical Properties Contiguity (C) Mean free path ( μ m) Ceramic Particle Size ( μ m) Ceramic Volume (% vol) Binder Volume (% vol) Porous Volume (% vol) Sample H v (1.0) [GPa] K ic [MPam 1/ 2 ]* Transverse Rupture Strength [MPa] Young’s Modulus [GPa] 0,36 1,09 2,56 80 19 1 SPS114,1 6,2/5,4/ 5,0 130,6 139,7 0,25 1,32 1,10 78 21 1 SPS214,4 8,5/8,0/ 7,9 448,6 324,1 0,28 1,27 1,95 80 19 1 SPS314,6 7,4/7,0/ 7,1 546,0 402,7 0,32 1,20 2,28 85 14 1 SPS416,0 6,7/5,1/ 6,0 522,3 318,4 * Ref, [22,23], [24,25]. Fig. 3. Characteristic SEM micrographs of microindentation marks carried out on polished surface of the sintered samples. J.M. C´ ordoba Gallego International Journal of Refractory Metals and Hard Materials 128 (2025) 107094 3 strengthening. From the slope of the lineal fitting of the data, the hardening of the material as a function of time is presented for every studied sample to every load in Fig. 5. At higher load lower hardening rate was observed. At higher load, the area under the indentation tip is bigger and the number of dislocations, which contributes to the strengthening, per unit area was lower than when low load was used. Fracture toughness has been evaluated for each sample using Eqs. (1)–(3). The results obtained are shown in Table 3, where slight differences in the calculated K ic can be seen depending on the expression used. (Shetty et al.) [22,23] KIC(P) = 0,0319×(P a× ∑l √)(1) (Shubert et al) [24] KIC(S) = 0,0028× Hv×P×∑l √(2) (Evans and Charles et al.) [25] KIC(EC) = 0,16×Hv× a √ (c a) 3 / 2 (3) where K ic [M Pa m 1/2 ], P: load [N], l: crack length: [mm], H v : hardness [N/mm 2 ], a: half-diagonal [mm] and c: crack+half diagonal [mm]. It should be noted at this point that the equation of Schubert et al. [24] was derived from that of Shetty et al. [22,23], and can be deduced by adjusting the units and experimental parameters. Even so, in our case it was used as a way of introducing different experimental parameters into an equation to obtain K IC, and be able to compare the data obtained with different experimental criteria. Finally, in this way, for K IC comparison, while the Evans and Charles equation, which was formed in the previous decade, was produced for materials in general, the Shetty and Shubert equations were constructed for composite materials, such as cemented carbides (WC-Co), similar to the material presented in this work. The average K ic values obtained from the equations used can be summarized as; SPS1, 5.6; SPS2, 8.1; SPS3, 7.1 and SPS4, 5.9. These values are closely related to the average grain size and the mean free path. The smaller the grain size or higher the mean free path is, the greater the fracture toughness. The volume percentage of grain borders or interfaces grows significantly with decreasing grain size or increasing mean free path. In materials such as intermetallic compounds and ceramics, which are usually extremely strong but very brittle, reducing the grain size or increasing the mean free path increment promotes ductility by increasing the probability of grain boundary sliding. Performance gains and a large processing advantage are provided by this. The TRS and Young’s modulus determined at room temperature of the studied samples are presented in Table 3. At constant sintering temperature, the tougher is the sintering conditions (higher time or higher applied pressure), the greater is the TRS. Samples SPS3 and SPS4 showed the highest combination of TRS and Young’s modulus in comparison with samples SPS1 and SPS2, showing the sample SPS1 the lowest values for all studied samples. The remarkable increment in TRS and Young’s modulus from sample SPS1 to samples SPS2, SPS3, and SPS4 can be explained in terms of contiguity and mean free path. The higher the mean free path and lower Fig. 4. Microhardness versus time of the indentation run showing the increment in the microhardness of the sample by working hardening. Fig. 5. Hardening rate (GPa/s) versus indentation load. J.M. C´ ordoba Gallego International Journal of Refractory Metals and Hard Materials 128 (2025) 107094 4 the contiguity is, the greater the TRS and Young’s modulus are [26,27]. The combination of values of ceramic particle size, contiguity and mean free path made of the sample SPS3 is the best match to have the highest TRS value. Characteristics SEM image of a fracture surface of sample SPS3 is presented in Fig. 6. In Fig. 6(a), a secondary electrons image shows the topography of the fracture surface, and in Fig. 6(b), a backscattering electron image shows the phase distribution associated with the different contrasts. The fracture surface observed of studied cermets showed mainly a typical trans-granular fracture. This kind of fracture mechanism showed the high adhesion between the ceramic phase and the intermetallic phase. It is suggested that tantalum presence influenced the interfacial energies of the system, yielding a stronger hard phase scaffolding. The cermet fracture surfaces become finer and the ratio of intergranular fracture increases to some extent with the increasing of finer ceramic grins content (Fig. 6, right top). It is evident that the metal binder is still covering the grains on the fracture surface, displaying dimples. The following describes the toughening mechanisms: the presence of coreless grains (samples SPS2 and SPS3) can be thought to be advantageous for the cermets’ toughness. It is simple to promote more fracture formation at the interface between the core and rim since the bonding force there is weaker than that of the homogeneous grain (samples SPS1 and SPS4). It implies that higher fracture toughness is achieved by cermets with a single phase and no core/rim interaction [28,29]. In comparison with other reports [30] where tantalum presence worsens the mechanical properties, the spark plasma sintering technique has proven to form an adequate environment for the obtaining of sintered bodies with outstanding mechanical properties. 4. Conclusions The effects of SPS parameters on microstructure and mechanical properties of Ti 0.9 Ta 0.1 C 0.5 N 0.5 /Co-based cermets were investigated and the conclusions obtained in this study can be summarized as follows: (i) Highly densified products were obtained independently of the conditions used. (ii) The best combination of mechanical properties and homogeneous microstructure was obtained when sintering time was held for 8 min at 30 MPa and 1150 ◦C. (iii) Hardening rate was quantified as a function of the microhardness versus run time. Considering the simplicity of the compositions, these cermets showed great potential for future applications. The presence of tantalum is thought to have a positive effect on the grain and/or phase borders. It is proposed that tantalum affects the system’s interfacial energies, resulting in the production of an intermetallic binder and a stronger hard phase skeleton. CRediT authorship contribution statement Jos´ e Manuel C´ ordoba Gallego: Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Software, Resources, Project administration, Methodology, Investigation, Formal analysis, Data curation, Conceptualization. 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. Acknowledgments I would like to thank Antonia Gallego C´ ardenas and Jos´ e Francisco C´ ordoba Sosa for their unwavering strength and support, for always pushing me to move forward even in the most difficult and hazardous situations. Data availability No data was used for the research described in the article. References [1] H. Yu, Y. Liu, Y. Jin, J. Ye, Effect of secondary carbides addition on the microstructure and mechanical properties of (Ti, W, Mo, V)(C, N)-based cermets, Int. J. Refrac. Metals Hard Mater. 29 (5) (2011) 586–590. [2] Dizaji V. Rahimi, , et al.J. Akbari, Microstructure and cutting performance investigation of Ti(C, N)-based cermets containing various types of secondary carbides, Int. J. Mach. Tools Manuf. 47 (5) (2007) 768–772. [3] P. Wu, Y. Zheng, Y. Zhao, H. Yu, Effect of TaC addition on the microstructures and mechanical properties of Ti(C,N)-based cermets, Mater. Des. 31 (7) (2010) 3537–3541. [4] P. Ettmayer, H. Kolaska, W. Lengauer, K. Dreyer, Ti(C,N) cermets-metallurgy and properties, Int. J. Refract. Met. Hard Mater. 13 (6) (1995) 343–351. [5] F. Qi, S. Kang, Study on microstructural changes in Ti(C, N)-NbC-Ni cermets, Mater. Sci. Eng. A 251 (1–2) (1998) 276–285. [6] Teruuchi K., Yano N., Titanium carbonitride-based cermet cutting insert, US Patent No. 5,518,822, 1996. [7] V.I. Tret’yakov, V.I. Mashevskaya, Effect of tantalum on the properties and structure of hard alloys based on titanium carbonitride, V.I, Powder Metal. Met. Ceram. 38 (1–2) (1999) 64–67. [8] Rudy E., Spinodal carbonitride alloys for tool and wear applications, US Patent No. 3,971,656, 1976. [9] J.M. C´ ordoba, M.J. Sayagu´ es, M.D. Alcal´ a, F.J. Gotor, Monophasic nanostructured powders of niobium, tantalum, and hafnium carbonitrides synthesized by a mechanically induced self-propagating reaction, J. Am. Ceram. Soc. 90 (2) (2007) 381–387. Fig. 6. SEM images of sample SPS3, (a) Secondary electron SEM image showing fracture surface topography, and (b) Backscattering SEM image showing the fracture surface phases composition distribution (dark grey, ceramic; bright grey, binder). J.M. C´ ordoba Gallego International Journal of Refractory Metals and Hard Materials 128 (2025) 107094 5 [10] J.M. C´ ordoba, M.J. Sayagu´ es, M.D. Alcal´ a, F.J. Gotor, Monophasic Ti y Nb 1-y C x N 1-x nanopowders obtained at room temperature by MSR, J. Mater. Chem. 17, [7] (2007) 650–653. [11] J.M. C´ ordoba, M.D. Alcal´ a, M.A. Avil´ es, M.J. Sayagu´ es, F.J. Gotor, New production of TiC x N 1-x -based cermets by one step mechanically induced self-sustaining reaction: powder synthesis and pressureless sintering, J. Eur. Ceram. Soc. 28 (10) (2008) 2085–2098. [12] J.M. C´ ordoba, M.A. Avil´ es, M.J. Sayagu´ es, M.D. Alcal´ a, F.J. Gotor, Synthesis of complex carbonitride powders Ti y M T1-y C x N 1-x (M T : Zr, V, Ta, Hf) via a mechanically induced self-sustaining reaction, J. Alloys Compd. 482 (1–2) (2009) 349–355. [13] Z.A. Munir, U. Anselmi-Tamburini, M. Ohyanagi, The effect of electric field and pressure on the synthesis and consolidation of materials: a review of the spark plasma sintering method, J. Mater. Sci. 41 (2006) 763–777. [14] M. Omori, Sintering, consolidation, reaction and crystal growth by the spark plasma system (SPS), Mater. Sci. Eng. A 287 (2) (2000) 183–188. [15] E. Khaleghi, Y.S. Lin, M.A. Meyers, E.A. Olevsky, Spark plasma sintering of tantalum carbide, Scripta Mater. 63 (2010) 577–580. [16] V. Mamedov, Spark plasma sintering as advanced PM sintering method, Powder Metall. 45 (4) (2002) 322–328. [17] C.C. Jia, H. Tang, X.Z. Mei, F.-Z. Yin, X.H. Qu, Spark plasma sintering on nanometer scale WC-Co powder, Mater. Letters 59 (19–20) (2005) 2566–2569. [18] J.M. C´ ordoba, et al., Spark plasma sintering of Ti x Ta 1-x C 0.5 N 0.5 -based cermets: effects of processing conditions on chemistry, microstructure and mechanical properties, Chem. Eng. J. 230 (2013) 558–566. [19] J. Rodríguez-Carvajal, R. Thiery, FullProf ILL. https://www.ill.eu/sites/fullprof/, May, 2010. [20] J.M. C´ ordoba, E. Chicardi, F.J. Gotor, Development of multicomponent-multiphase materials based on (Ti,Ta,Nb)C x N 1-x carbonitride solid solutions, Chem. Eng. J. 192 (2012) 58–661. [21] U.A. Tamburini, S. Gennari, J.E. Garay, Z.A. Munir, Fundamental investigations on the spark plasma sintering/synthesis process II. Modeleing of current and temperature distributions, Mater. Sci. Eng. A 394 (2005) 139–148. [22] D.K. Shetty, I.G. Wright, P.N. Mincer, A.H. Clauer, Indentation fracture of. WC-Co cermets, J. Mater. Sci. 20 (1985) 1873–1882. [23] D.K. Shetty, I.G. Wright, On estimating fracture toughness of cemented carbides from Palmqvist crack sizes, J. Mater. Sci. Lett. 5 (1986) 365–368. [24] W.D. Schubert, H. Neumeister, G. Kinger, Hardness to toughness relationship of fine-grained WC-Co hardmetals, Int. J. Refract. 16 (1998) 133–142. [25] A.G. Evans, E.A. Charles, Fracture toughness determination by. indentation, J. Am. Ceram. Soc. 59 (1976) 371–372. [26] K. Chang-Soo, Microstructural-mechanical property relationships in tungsten carbide-cobalt composites, Doctoral Thesis, Carnegie Mellon University, 2004. Publication N◦: AAI3146413; ISBN: 9780496075584; Advisor: Massa, Ted R, htt p://mimp.materials.cmu.edu/~gr20/theses/Kim_thesis.pdf. [27] Z. Guo, J. Xiong, M. Yang, C. Jiang, WC–TiC–Ni cemented carbide with enhanced properties, J. Alloys Compd. 465 (2008) 157–162. [28] A. Celli, A. Tucci, L. Esposito, C. Palmonari, Fracture analysis of cracks in alumina–zirconia composites, J. Eur. Ceram. Soc. 23 (2003) 469–479. [29] S. Park, S. Kang, Toughened ultra-fine (Ti, W)(CN)–Ni cermets, Scr. Mater. 52 (2005) 129–133. [30] M. Fides, et al., TiTaCN-Co cermets prepared by mechanochemical technique: microstructure and mechanical properties, Procedia Eng. 149 (2016) 87–93. J.M. C´ ordoba Gallego International Journal of Refractory Metals and Hard Materials 128 (2025) 107094 6