Microstructure and compression behavior of Ag–W metal matrix composite produced from core–shell powder by spark plasma sintering: case study

Vol.:(0123456789) The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 https://doi.org/10.1007/s00170-025-15944-7 ORIGINAL ARTICLE Microstructure andcompression behavior ofAg–W metal matrix composite produced fromcore–shell powder byspark plasma sintering: case study AngelinaStrakošová1· DrahomírDvorský1· FilipPrůša2· OrsolyaMolnárová1· StanislavHabr1· JakubSvoboda3· IvonaSedlářová2· DaliborVojtěch2· PavelLejček1 Received: 19 March 2025 / Accepted: 5 June 2025 / Published online: 24 June 2025 © The Author(s) 2025 Abstract Metal matrix composites represent an interesting class of materials with an exclusive combination of properties. In this study, a unique Ag–W metastable metal matrix composite was produced from W@Ag core–shell powders using a spark plasma sintering technique at a temperature of 700°C and a pressure of 80MPa. The microstructures of a default powder and asproduced composite were observed by scanning electron microscopy equipped with energy-dispersive X-ray spectroscopy and electron backscatter diffraction. As expected, the composite is characterized by a dual microstructure: a soft matrix of pure Ag with a submicrometer grain size reinforced by W particulates with a grain size of up to 30μm. In addition, tensile and compression tests were performed with a deformation rate of 10–3 s−1 at ambient temperature. The value of the compression yield stress of Ag–W MMC is higher than the compression yield stress of pure Ag by approximately 467%. Observation of the microstructure of the deformed composite material revealed that the interface between the matrix and the reinforced particles is the weakest place, which is a key factor influencing the performance and properties of the composite material. Keywords Core–shell powder· Silver· Tungsten· Spark plasma sintering· Metal matrix composite· Compression deformation 1 Introduction Metal matrix composites (MMCs) represent a class of highly sophisticated materials with a unique combination of properties, including, e.g., enhanced strength, wear resistance, and reduced thermal expansion [1, 2], that are well-suited for a diverse array of applications such as aerospace, military, sports, and electronics [3, 4]. MMCs include a continuous metallic matrix and a meticulously selected reinforcement, which is deliberately chosen based on its specific properties and the desired features of the composite material [3, 5, 6]. It is known that the shape of reinforcement (fibers and particles) affects the structure and properties of composite materials [7]. A trendy branch of MMC preparation represents the methods of powder metallurgy (PM), such as hot extrusion and hot pressing [8]. Besides them, spark plasma sintering (SPS) is a promising method in this respect if small parts should be prepared. Compared to other PM methods, SPS allows the preservation of fine microstructure [9, 10] due to the short-term nature of the process, which typically combines only a few minutes of sintering and consolidation into a single operation step [8, 11]. A wide range of materials can be produced by the SPS technique [12, 13], for example, high-entropy alloys [9], composites with extremely high strength [14], and intermetallics [15]. Recently, core–shell powders have become a popular material used for MMC production [16–20]. A combination of the SPS technique with the usage of core–shell powders enables the production of metastable metal matrix composites (m-MMCs) thanks to the sintering of the shell surfaces and suppression of the diffusion processes between the core and shell [21]. Recently, we studied the structure and plastic * Angelina Strakošová strakoso[email protected] 1 Institute ofPhysics, Czech Academy ofSciences, Na Slovance 2, 18200Prague, CzechRepublic 2 University ofChemistry andTechnology, Technická 5, 16628Prague, CzechRepublic 3 SAFINA a.s, Vídeňská 104, 25250Vestec, CzechRepublic 1572 The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 deformation of Ag–Cu m-MMCs prepared from Cu@Ag core–shell powders, where the core and shell materials show very similar mechanical properties [21, 22]. On the other hand, it could be a very interesting topic to combine two materials with diametrically different properties to produce MMCs. One of such materials could be MMCs, which are used in electronics as electrical contacts and are a combination of two immiscible metals, such as silver and tungsten [23]. Such composites, which can also have antibacterial potential, have usually been formed by spherical particulates in a soft matrix and are usually produced by hot pressing Ag and W powders [23, 24]. Only a few attempts have been performed to produce micro-sized MMCs from Ag@ (W/WO3) core–shell powders [16, 25–27] and study their mechanical and electric properties [28–30]. Unfortunately, the present knowledge on Ag-based MMCs is scarce. Therefore, their production and examination of the microstructure and other behaviors, such as the mechanical properties of Ag-W MMCs, are intriguing and important topics. In the present work, we describe the microstructure and mechanical behavior of a model MMC combining a very soft Ag matrix (density = 10.5 g cm−3, hardness = 300–700 MPa; bulk resistivity = 1.63 µΩ cm [31]) reinforced with high-strength W particulates (density = 19.3 g cm−3; hardness = 1200–4000 MPa; bulk resistivity = 5.44 µΩ cm [31]) prepared from W@Ag core–shell powder using the SPS technique. Compared to the abovementioned Ag–Cu system, Ag–W MMC reported here exhibits entirely different compression behavior, not only in terms of mechanical properties but also due to its complete immiscibility. 2 Experimental 2.1 Starting material The W@Ag core–shell powder with a volume ratio of W and Ag 50:50 was manufactured in SAFINA a.s. The production of the powder consists of two steps: i) gas atomization of W powder and ii) dosage of AgNO3 in an alkaline solution for obtaining pure Ag for precipitation of Ag on W particulates. The powder particle distribution was measured by laser diffraction on the Malvern Panalytical Mastersizer 3000 machine (Malvern, United Kingdom). 2.2 Production ofAg–W MMC The core–shell powder was compacted by SPS technology using FCT Systeme HP-D 10 device (Rauenstein, Germany) to produce a rounded sample with a diameter of 20 mm and a height of 6mm. For each sample, 30 g of powder was used. The parameters of the compaction process were identical to those described in detail for the production of Cu–Ag m-MMCs in the work [21], i.e., i) a force of 25 kN (corresponds to a pressure of 80 MPa), ii) a heating rate of 100 °C min−1, iii) a compaction temperature of 700 °C, iv) a dwell time of 5min, and v) a cooling segment—turning off the heating to achieve the maximal cooling rate of the device. The time dependencies of the recorded height reduction, step temperature, current flow, and applied force during the process are shown in Fig.1. 2.3 MMC characterization A morphology of the default W@Ag core–shell powder, as well as its microstructure and the microstructure of the Ag–W MMC, were characterized using scanning electron microscopy (SEM, Tescan Mira, Czech Republic) equipped by energy-dispersive X-ray spectroscopy (EDS, Oxford Instruments X-Max 20, United Kingdom). The detailed microstructural characterization of the as-produced Ag–W MMC was performed on a 50 × 80 µm area with a step size of 0.1 µm using SEM (FEI Quanta 3D FEG, Thermo Fisher, Czech Republic), including electron backscatter diffraction (EBSD, Gatan, The United States of America). The obtained EBSD data were processed using TSL OIM8 software (v. 8.0). All the microstructures were observed on the ground and polished surfaces. 2.4 Deformation tests Tensile and compression tests were performed to investigate the mechanical properties of the obtained Ag–W MMC. In both cases, the Instron 5882 machine (Instron, The United States of America) with a deformation rate of 10–3 s−1 at ambient temperature was used. Tensile tests were performed on samples with a “dog-bone” shape with dimensions of 7.0-mm length, 2.0-mm width, and 0.55-mm thickness. Cuboid samples with a 6.0-mm height and a square base of 4 × 4mm were used for compression tests. Samples for tensile and compression tests were prepared by electro-erosive Fig. 1 Recorded height reduction (blue), applied force (black), current flow (gray), and sample temperature (red) during the SPS process 1573The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 cutting from the samples produced by SPS. Tensile/compression true stress–true strain curves are automatically determined by Bluehill software, which is used by the Instron 5882 machine. The cross-sections of the deformed materials were observed using SEM Tescan Mira. 3 Results anddiscussion 3.1 Characterization ofthestarting material The distribution of particle size in the W@Ag core–shell powder is shown in Fig.2. The particle sizes exhibit a typical Gaussian distribution with dimensions ranging from 20 to 80 µm with a median value of 37.1 µm, which is approximately 10 µm larger than the size of Cu@Ag core–shell described in [21]. The morphology of the W@Ag core–shell powder is characterized by a spherical/rounded shape with the surface looking like the orange peel texture (Fig.3). There can also be seen voids on the surface of some particles (Fig.3a, yellow arrows). The presence of these defects (as well as the surface roughness described in the study [21]) can affect the microstructure of the composite produced by the SPS process. If the void is too deep and reaches the W core, islands of sintered W can be observed within the microstructure of the produced composite. The microstructure and the chemical composition of the W@Ag core–shell powder cross-section are shown in Fig.4. Unfortunately, a height difference between the hard W core and the very ductile Ag shell occurred during the metallographic preparation of the sample due to a different response of Ag and W to the grinding and subsequent polishing. In addition, wide hollow pores are visible between the two component materials, an example of which is marked by a red arrow in Fig.4. The origin of this defect is unclear; it could occur either during the metallographic preparation or the manufacturing method during the Ag deposition onto the W core. In this case, it could result from the 25% difference in lattice parameters of fcc Ag (a = 0.4086 nm [32, 33]) and bcc W (a = 0.3155 nm [34, 35]). Fig. 2 Size distribution and cumulative curve of the W@Ag core– shell powder Fig. 3 SEM micrographs of the W@Ag core–shell powder morphology with different magnifications. Yellow arrows in (a) mark the voids in the shell of the particles 1574 The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 3.2 Characterization oftheAg–W MMC The microstructure of the Ag–W MMC is shown in Fig.5. It is seen that the SPS process made it possible to produce the MMC characterized by Ag matrix and W particulates. However, a lot of W clusters are observed here. As was mentioned above, it can be a result of a defective surface of the core–shell powder particles (Fig.3a, yellow arrows), which makes it possible to join the W core together during the SPS process. Figure5 also shows a high number of voids in the MMC microstructure. It is apparent that the shape and size of the voids correspond to the size and shape of W particulates. These voids occurred because of dropping out of the W particulates during the grinding and polishing of the material. The dropping out of the W particulates is a consequence of the fact that the W and Ag do not show any solubility according to the phase diagram [36]. A weak bonding between Ag and W was also observed in [16], where the W–Ag composites were prepared with different Ag contents. As a result, it was impossible to characterize the porosity of the Ag–W MMC. We might expect that the true porosity of the Fig. 4 SEM micrograph and EDS elements distribution maps of W and Ag at the cross-section of the W@Ag core–shell powder. A red arrow shows the area of bad adhesion between the shell and core Fig. 5 SEM micrograph of the Ag–W MMC produced by SPS, different magnifications 1575The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 compacted MMC will be on the same level as that determined for Ag–Cu MMCs, i.e., of the order of 10−2–10−1%. Figure6 depicts the EBSD results of the Ag–W MMC produced by the SPS technique. Inverse pole figure (IPF) maps show a dual structure of the MMC—large W particulates consisting of a few grains while being embedded in the Ag matrix with nanosized grains. The grain size analysis of both W and Ag (Fig.7) documents that W grains have sizes in a range from a few micrometers up to almost 30 µm. Conversely, the majority of Ag grain sizes range between 500 nm and 1µm. The difference in grain sizes of the two components of the MMC results from the production way of the core–shell powder— atomization of W particles and deposition of the Ag shell. EDS element distribution maps (Fig.6) show some fine particles of W in the Ag matrix. It is apparent that their distribution, size, and shape are completely coincidental. Figure5b shows the same type of pure elemental particles of W in the Ag matrix. Therefore, their occurrence can be attributed to the metallographic sample preparation when the hard but brittle W is broken and stuffed into a very soft and ductile Ag. 3.3 Mechanical testing To investigate the mechanical properties of the Ag–W MMC produced by SPS technique, tensile and compression tests were aimed to be performed. Regrettably, it was impossible to obtain any reasonable result from tensile tests. Samples of a “dog-bone” shape tend to break apart during their insertion into the device clamp. For illustration, Fig. 6 IPF and EDS maps of the Ag–W MMC obtained from EBSD analysis Fig. 7 The plots of the grain sizes in the Ag–W m-MMC: (a) W and (b) Ag 1576 The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 the fracture surface of such a damaged sample is shown in Fig.8. The fracture surface shows nondeformed W particulates. Moreover, when separated from the Ag matrix, these particulates left the dimples with corresponding shapes and sizes that are almost 100% smooth. It indicates that there was the weakest cohesion between the matrix and reinforcement particulates. This finding can be related to the pulling out of the W particulates during grinding and polishing (Fig.5). Figure8 also shows that the area around the spherical particulates exhibits a fracture, characterized by dimples, confirming a ductile behavior of the Ag matrix. In fact, this fracture runs along the weakest path of the matrix among close particulates where the lowest force is necessary to separate the material. The EDS element distribution maps (Fig.8) prove the presence of nondeformed W particulates in the ductile Ag matrix. Similar behavior of composite material failure by a decohesion along the interface between the hard particulates (ZrO2/ SiO2) and the soft matrix (Al) was described in the study [7]. In addition, the authors of [37] described three types of particulate defects at the fracture surface of 6082Al matrix composite reinforced with Mo particulates: (i) fractured particulates, (ii) pull-out particulates, and (iii) pull-out particulates pits. In the present study, the fractured particulates were not observed; however, the (ii) and (iii) types are shown in Fig.8 and marked by blue and red arrows, respectively. The dependence of the true stress–true strain in compression of Ag–W MMC, together with those of pure Ag as well as of recently studied Ag–Cu MMC [21], is shown in Fig.9. As expected, the value of the compression yield stress (CYS) of Ag–W MMC is much higher than the CYS of pure Ag. On the other hand, it is nearly the same as that of Ag–Cu MMC, containing much softer reinforcing particulates. According to the recent work [21], the CYS of Ag–Cu MMC (153 ± 10 MPa) is higher by 467% than that of the pure Ag matrix (27 ± 6 MPa). The yield stress, 𝜎C , of a compressed composite can be written as follows [21, 38–41]: (1) 𝜎C=𝜎m+f(Δ𝜎D,Δ𝜎LT ,Δ𝜎GB,Δ𝜎Or,Δ𝜎SS,…) where 𝜎m is the CYS of the matrix, Δ𝜎D is the stress of the dislocations, Δ𝜎TM is a thermal mismatch, Δ𝜎GND is the geometrical necessity, Δ𝜎LT is the contribution of the load transfer, Δ𝜎GB is the contribution of grain boundary (Hall–Petch) strengthening, Δ𝜎Or is the contribution of the particle (Orowan) strengthening, and Δ𝜎SS is the stress contribution from the solid solution. The function in Eq. (1) is complex [42, 43]; however, it can also be considered as a simple sum of the components [38, 39, 44]. In the Ag–W MMCs, we can exclude the contribution of Δ𝜎SS as Ag and W are immiscible. The contribution Δ𝜎D may be written as [45]: with the constant A, the shear modulus G, the Burgers vector b, and the summary dislocation density 𝜌dis . Accepting A = 0.55 [45], GAg = 24.3 GPa [46] and |b(Ag)|= 0.290 (2) Δ 𝜎 D =AG � 𝐛 �√ 𝜌 dis Fig. 8 SEM micrograph and EDS maps of the Ag–W m-MMC fracture surface after the “tensile” test. Blue arrows depict the W particulates, and red arrows mark the holes after dropped W particulates Fig. 9 Compression true stress–true strain curves of the Ag–W MMC. For comparison, the compression curves of pure Ag and Ag– Cu (45:55) MMC [21] are also shown 1577The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 nm [47], and supposing the dislocation density in Ag to be 8 × 1013 m–2, i.e., similar to that measured in the Ag–Cu MMC [22], the mean contribution of dislocations in the matrix to the stress is Δ𝜎D = 35 MPa. The form of the term Δ𝜎LT can be given as follows [48]: where 𝜅 is the aspect ratio of the reinforcements, 𝜎y is the yield stress of the particulate, and fv is the volume fraction of the reinforcement phase. To estimate its contribution, we can use the values of 𝜅 = 1, fv = 0.50, and 𝜎y = 56 MPa [21]. Then, Δ𝜎LT = 14 MPa. The contribution of the grain boundaries (Hall–Petch contribution) in polycrystalline materials with the grain size dg causes an increase in the stress [49]: with KGB being constant. For Ag, we can use the values of dg = 0.75 μm and KGB = 2.51 MPa mm1/2 [49]. Supposing that the active area of Ag is 50% in the MMC, which is composed of 50% of both Ag and W, we obtain the value of Δ𝜎GB = 45 MPa. Then, the effect of the Orowan strengthening must not be considered here. The large non-deformable particulates serve in the MMC as large objections reducing the active area. This effect was considered in the Hall–Petch strengthening above. Supposing a simple addition of individual contributions, the summary increase of 𝜎m in the composite is then of about 94 MPa. This estimated value is close to the difference between the measured values of CYS of the Ag–W MMC (135 MPa) and of the Ag matrix (35 MPa) [21], i.e. 100 MPa. The fact that the values of CYS of Ag–W and Ag–Cu MMCs are very similar (Fig.9) results from the fact that plastic deformation starts when the dislocations become mobile in the Ag matrix and that the modifications of the values of CYS in both cases are dominated by the behavior of the Ag matrix. In contrast to pure Ag and Ag–Cu MMC exhibiting a similar course of the true stress–true strain curve, its course for Ag–W MMC is quantitatively different during compression (Fig.9). This is caused by the fact that W particulates do not deform plastically, thus limiting plastic deformation to a more limited volume than it is in the case of both pure Ag and Ag–Cu MMC. In the latter one, plastic deformation also occurs in the particulates, which strongly contributes to an increase of the strengthening and suppresses the fracture of the material. In this respect, the W particulates represent an objection to the plastic deformation of the MMC. Supposing the value of Young’s modulus of W to be 370 (3) Δ 𝜎LT = 1 2 𝜎y𝜅fv , (4) Δ 𝜎GB = K GB √ d g , GPa [50], the elastic deformation at 210 MPa, where the plastic deformation of the MMC is 15.4% (Fig.9), is about 0.068% only. This causes large internal stress at the Ag/W interface, which results eventually in fracturing the MMC. Additionally, a rather flat course of the deformation curve can be related to a relatively large size of the reinforcing particulates [51]. To characterize the plasticity during the compression, we can evaluate the work hardening rate as the change of true stress, σ, with true strain, ε, under constant deformation rate, 𝜀 [52], and the work hardening exponent n [52], The dependences of both parameters on the true strain are shown in Fig.10. Due to the flat dependence of σ vs. ε discussed above, the work hardening rate is quickly decreasing to zero at the ultimate compression stress (UCS) and reaching negative values during the fracturing of the sample along the Ag/W interfaces as well as in the Ag matrix (Fig.10a). It contrasts with the behavior of pure Ag and Ag–Cu MMC characterized by a continuous increase of the stress with the strain (Fig.9). This difference results from the fact that W particulates do not deform plastically while the Cu ones in Ag–Cu MMCs do, and in addition, Cu particulates in the corresponding MMC as well as pure Ag-contained large grains [21] which strengthen similarly as documented in Fig.10a. The work hardening exponent narrates about the mechanism of plastic deformation. The values of n = 0.15–0.18 found for the Ag–W MMC (Fig.10b) indicate that the deformation is realized by the slip and climb of dislocations in fine-grained Ag. A similar mechanism of deformation was reported for fine-grained Ti [53]. A similar value of n = 0.21 was also measured at the early stage of the deformation of Ag–Cu MMC, referring to the starting of the dislocation motion in the fine-grained Ag matrix. However, with increasing stress, deformation also occurs in coarsegrained Cu particulates, which is reflected in an increased value of n (n = 0.31). As an increase of n is attributed to deformation twinning in Ti [53], we may deduce that twinning is also operating in the Cu particulates. In the case of Ag possessing a grain size of 0.5 mm, twinning can easily occur already at the early stage of deformation and lead to the subsequent increase of n value from 0.28 to 0.53. In Ag–W MMC, twinning was suppressed due to small grain size and non-deformability of W particles; thus, an insignificant increase of n was observed. (5) Θ=(𝜕𝜎 𝜕𝜀 )𝜀 , (6) n =( 𝜕ln 𝜎 𝜕ln 𝜀 ) 𝜀 1578 The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 The microstructure of the fracture surfaces of Ag–W MMC after compression deformation is shown in Fig.11. It is apparent that the W particulates remain unchanged (in terms of size and shape) after compression tests, which confirms that W was deformed elastically only. The abovementioned concentration of the stress at the Ag/W interfaces results in the occurrence of cracking there. Then, the fracture continues in a ductile manner through the joining Ag matrix. There are two types of morphology of the fracture: the dimple character of simple fracture and the shear in the Ag matrix caused by the hard particulates moving during compression (Fig.11). Nevertheless, the deformability of the Ag–W MMC, exhibiting more than 15%, is relatively very good. Fig. 10 Characteristics of plastic deformation of Ag–W MMC. (a) work hardening rate; (b) work hardening exponent. The dependences of pure Ag and of Ag–Cu MMC are shown for comparison Fig. 11 SEM micrographs of the fracture surface of the Ag–W MMC after compression tests at different magnifications 1579The International Journal of Advanced Manufacturing Technology (2025) 139:1571–1580 4 Conclusions In this study, an Ag–W MMC was successfully fabricated using SPS from W@Ag core–shell powder. As a result, a dual structure characterized by an Ag matrix with grain size ranging from 500 nm to 1μm reinforced by W particulates with grain size of up to 30 μm was produced. The increase in CYS of the MMC compared to that of pure Ag by approximately 467% is attributed to the strengthening effects of dislocation (by 35 MPa), load transfer (by 14 MPa), grain boundaries (Hall–Petch, by 45 MPa), but also to the strengthening of thermal mismatch and geometrical necessity, as evaluated quantitatively. A low and nearly constant value of the work hardening exponent (n = 0.15–0.18) implies that plastic deformation occurs in the Ag–W MMC via dislocation motion in the matrix, while the particulates do not deform during compression. The interface between the matrix and reinforcement particulates, however, was identified as the weakest link in the MMC structure. It suggests that further research is needed, as different conditions of the SPS process or using other methods of PM, to improve the strength of this interface. Author contribution All authors made significant contributions to the study conception and design. Material preparation, data collection, and analysis were performed by Angelina Strakošová, Drahomír Dvorský, Filip Průša, Orsolya Molnárová, Stanislav Habr, Jakub Svoboda, and Ivona Sedlářová. The first draft of the manuscript was written by Angelina Strakošová, Drahomír Dvorský, and Pavel Lejček, and all authors commented on previous versions of the manuscript. Funding was secured by Pavel Lejček. All authors read and approved the final manuscript. Funding Open access publishing supported by the institutions participating in the CzechELib Transformative Agreement. This work was financially supported by the Czech Science Foundation under the grant No. 23-05139S. AS acknowledges the grant of the Czech Academy of Sciences under grant No. L100102403. FP and DV acknowledge financial support from the project “Mechanical Engineering of Biological and Bio-inspired Systems,” funded as project No. CZ.02.01.01/00/22_008/0004634 by Programme Johannes Amos Commenius, call Excellent Research. We acknowledge CzechNanoLab Research Infrastructure supported by MEYS CR (LM2023051). Grantová Agentura České Republiky, 23-05139S, Pavel Lejček;Akademie Věd České Republiky,L100102403,Angelina Strakošová;Programme Johannes Amos Commenius,call Excellent Research, CZ.02.01.01/00/22_008/0004634; Ministerstvo Školství,Mládeže a Tělovýchovy,LM2023051. Data availability The data used in this manuscript are available in the Zenodo repository at the following link:https:// doi. org/ 10. 5281/ zenodo. 14775 989. Declarations Conflict of interest The authors declare no competing interests. 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