CoCrFeNi is a well-studied face-centered cubic (fcc) high-entropy alloy (HEA) with excellent ductility but limited strength. The focus of this study is on improving the balance of strength and ductility of such HEAs by adding different amounts of SiC using the arc melting method. It has been established that the presence of chromium in the base HEA causes the decomposition of SiC during melting. Thus, the interaction of free carbon with chromium leads to the in situ formation of chromium carbides, while free silicon remains in solution in the base HEA and/or interacts with the elements that make up the base HEA to form silicides. As the SiC content increases, the microstructure phase changes in the following sequence: fcc → fcc + eutectic → fcc + chromium carbide flakes → fcc + chromium carbide flakes + silicide → fcc + chromium carbide flakes + silicide + graphite balls / graphite flakes. The resulting composites exhibit a very wide range of mechanical properties (yield strength ranging from 277 MPa at over 60% elongation to 2522 MPa at 6% elongation) compared to conventional alloys and high entropy alloys. Some of the high entropy composites developed show an excellent combination of mechanical properties (yield strength 1200 MPa, elongation 37%) and occupy previously unattainable regions on the yield stress-elongation diagram. In addition to remarkable elongation, the hardness and yield strength of HEA composites are in the same range as bulk metallic glasses. Therefore, it is believed that the development of high-entropy composites can help achieve an excellent combination of mechanical properties for advanced structural applications.
The development of high entropy alloys is a promising new concept in metallurgy1,2. High entropy alloys (HEA) have shown in a number of cases an excellent combination of physical and mechanical properties, including high thermal stability3,4 superplastic elongation5,6 fatigue resistance7,8 corrosion resistance9,10,11, excellent wear resistance12,13,14,15 and tribological properties15 ,16,17 even at high temperatures18,19,20,21,22 and mechanical properties at low temperatures23,24,25. The excellent combination of mechanical properties in HEA is usually attributed to four main effects, namely high configurational entropy26, strong lattice distortion27, slow diffusion28 and cocktail effect29. HEAs are usually classified as FCC, BCC and HCP types. FCC HEA typically contains transition elements such as Co, Cr, Fe, Ni and Mn and exhibits excellent ductility (even at low temperature25) but low strength. BCC HEA is usually composed of high density elements such as W, Mo, Nb, Ta, Ti and V and has very high strength but low ductility and low specific strength30.
The microstructural modification of HEA based on machining, thermomechanical processing and the addition of elements has been investigated to obtain the best combination of mechanical properties. CoCrFeMnNi FCC HEA is subjected to severe plastic deformation by high-pressure torsion, which leads to a significant increase in hardness (520 HV) and strength (1950 MPa), but the development of a nanocrystalline microstructure (~50 nm) makes the alloy brittle31. It has been found that the incorporation of twinning ductility (TWIP) and transformation induced plasticity (TRIP) into CoCrFeMnNi HEAs confers good work hardenability resulting in high tensile ductility, albeit at the expense of actual tensile strength values. Below (1124 MPa) 32. The formation of a layered microstructure (consisting of a thin deformed layer and an undeformed core) in the CoCrFeMnNi HEA using shot peening resulted in an increase in strength, but this improvement was limited to about 700 MPa33. In search of materials with the best combination of strength and ductility, the development of multiphase HEAs and eutectic HEAs using additions of non-isoatomic elements has also been investigated34,35,36,37,38,39,40,41. Indeed, it has been found that a finer distribution of hard and soft phases in eutectic high-entropy alloys can lead to a relatively better combination of strength and ductility35,38,42,43.
The CoCrFeNi system is a widely studied single-phase FCC high-entropy alloy. This system exhibits fast work hardening properties44 and excellent ductility45,46 at both low and high temperatures. Various attempts have been made to improve its relatively low strength (~300 MPa)47,48 including grain refinement25, heterogeneous microstructure49, precipitation50,51,52 and transformation-induced plasticity (TRIP)53. Grain refinement of cast face-centered cubic HEA CoCrFeNi by cold drawing under severe conditions increases the strength from about 300 MPa47.48 to 1.2 GPa25, but reduces the loss of ductility from more than 60% to 12.6%. The addition of Al to the HEA of CoCrFeNi resulted in the formation of a heterogeneous microstructure, which increased its yield strength to 786 MPa and its relative elongation to about 22%49. CoCrFeNi HEA was added with Ti and Al to form precipitates, thereby forming precipitation strengthening, increasing its yield strength to 645 MPa and elongation to 39%51. The TRIP mechanism (face-centered cubic → hexahedral martensitic transformation) and twinning increased the tensile strength of CoCrFeNi HEA to 841 MPa and elongation at break to 76%53.
Attempts have also been made to add ceramic reinforcement to the HEA face centered cubic matrix to develop high entropy composites that can exhibit a better combination of strength and ductility. Composites with high entropy have been processed by vacuum arc melting44, mechanical alloying45,46,47,48,52,53, spark plasma sintering46,51,52, vacuum hot pressing45, hot isostatic pressing47,48 and the development of additive manufacturing processes43,50. Carbides , oxides and nitrides such as WC44, 45, 46, Al2O347, SiC48, TiC43, 49, TiN50 and Y2O351 have been used as ceramic reinforcement in the development of HEA composites. Choosing the right HEA matrix and ceramic is especially important when designing and developing a strong and durable HEA composite. In this work, CoCrFeNi was chosen as the matrix material. Various amounts of SiC were added to the CoCrFeNi HEA and their effect on the microstructure, phase composition, and mechanical properties was studied.
High-purity metals Co, Cr, Fe, and Ni (99.95 wt %) and SiC powder (purity 99%, size -400 mesh) in the form of elementary particles were used as raw materials for the creation of HEA composites. The isoatomic composition of the CoCrFeNi HEA was first placed in a hemispherical water-cooled copper mold, and then the chamber was evacuated to 3·10-5 mbar. High purity argon gas is introduced to achieve the vacuum required for arc melting with non-consumable tungsten electrodes. The resulting ingots are inverted and remelted five times to ensure good homogeneity. High-entropy composites of various compositions were prepared by adding a certain amount of SiC to the resulting equiatomic CoCrFeNi buttons, which were re-homogenized by five-fold inversion and remelting in each case. The molded button from the resulting composite was cut using EDM for further testing and characterization. Samples for microstructural studies were prepared according to standard metallographic methods. First, the samples were examined using a light microscope (Leica Microscope DM6M) with the software Leica Image Analysis (LAS Phase Expert) for quantitative phase analysis. Three images taken in different areas with a total area of about 27,000 µm2 were selected for phase analysis. Further detailed microstructural studies, including chemical composition analysis and element distribution analysis, were carried out on a scanning electron microscope (JEOL JSM-6490LA) equipped with an energy dispersive spectroscopy (EDS) analysis system. The characterization of the crystal structure of the HEA composite was performed using an X-ray diffraction system (Bruker D2 phase shifter) using a CuKα source with a step size of 0.04°. The effect of microstructural changes on the mechanical properties of HEA composites was studied using Vickers microhardness tests and compression tests. For the hardness test, a load of 500 N is applied for 15 s using at least 10 indentations per specimen. Compression tests of HEA composites at room temperature were carried out on rectangular specimens (7 mm × 3 mm × 3 mm) on a Shimadzu 50KN universal testing machine (UTM) at an initial strain rate of 0.001/s.
High entropy composites, hereinafter referred to as samples S-1 to S-6, were prepared by adding 3%, 6%, 9%, 12%, 15%, and 17% SiC (all by weight%) to a CoCrFeNi matrix. respectively. The reference sample to which no SiC was added is hereinafter referred to as sample S-0. Optical micrographs of the developed HEA composites are shown in Figs. 1, where, due to the addition of various additives, the single-phase microstructure of the CoCrFeNi HEA was transformed into a microstructure consisting of many phases with different morphology, sizes, and distribution. The amount of SiC in the composition. The amount of each phase was determined from image analysis using LAS Phase Expert software. The inset to Figure 1 (upper right) shows an example area for this analysis, as well as the area fraction for each phase component.
Optical micrographs of the developed high-entropy composites: (a) C-1, (b) C-2, (c) C-3, (d) C-4, (e) C-5 and (f) C-6. The inset shows an example of contrast-based image phase analysis results using the LAS Phase Expert software.
As shown in fig. 1a, a eutectic microstructure formed between the matrix volumes of the C-1 composite, where the amount of the matrix and eutectic phases is estimated as 87.9 ± 0.47% and 12.1% ± 0.51%, respectively. In the composite (C-2) shown in Fig. 1b, there are no signs of a eutectic reaction during solidification, and a microstructure completely different from that of the C-1 composite is observed. The microstructure of the C-2 composite is relatively fine and consists of thin plates (carbides) uniformly distributed in the matrix phase (fcc). The volume fractions of the matrix and carbide are estimated at 72 ± 1.69% and 28 ± 1.69%, respectively. In addition to the matrix and carbide, a new phase (silicide) was found in the C-3 composite, as shown in Fig. 1c, where the volume fractions of such silicide, carbide, and matrix phases are estimated at about 26.5% ± 0.41%, 25.9 ± 0.53, and 47.6 ± 0.34, respectively. Another new phase (graphite) was also observed in the microstructure of the C-4 composite; a total of four phases were identified. The graphite phase has a distinct globular shape with dark contrast in optical images and is only present in small amounts (estimated volume fraction is only about 0.6 ± 0.30%). In composites C-5 and C-6, only three phases were identified, and the dark contrasting graphite phase in these composites appears in the form of flakes. Compared to the graphite flakes in Composite S-5, the graphite flakes in Composite S-6 are wider, shorter, and more regular. A corresponding increase in graphite content was also observed from 14.9 ± 0.85% in the C-5 composite to about 17.4 ± 0.55% in the C-6 composite.
To further investigate the detailed microstructure and chemical composition of each phase in the HEA composite, samples were examined using SEM, and EMF point analysis and chemical mapping were also performed. The results for composite C-1 are shown in fig. 2, where the presence of eutectic mixtures separating the regions of the main matrix phase is clearly seen. The chemical map of composite C-1 is shown in Fig. 2c, where it can be seen that Co, Fe, Ni, and Si are uniformly distributed in the matrix phase. However, a small amount of Cr was found in the matrix phase compared to other elements of the base HEA, suggesting that Cr diffused out of the matrix. The composition of the white eutectic phase in the SEM image is rich in chromium and carbon, indicating that it is chromium carbide. The absence of discrete SiC particles in the microstructure, combined with the observed low content of chromium in the matrix and the presence of eutectic mixtures containing chromium-rich phases, indicates the complete decomposition of SiC during melting. As a result of the decomposition of SiC, silicon dissolves in the matrix phase, and free carbon interacts with chromium to form chromium carbides. As can be seen, only carbon was qualitatively determined by the EMF method, and the phase formation was confirmed by the identification of characteristic carbide peaks in the X-ray diffraction patterns.
(a) SEM image of sample S-1, (b) enlarged image, (c) element map, (d) EMF results at indicated locations.
The analysis of composite C-2 is shown in fig. 3. Similar to the appearance in optical microscopy, SEM examination revealed a fine structure composed of only two phases, with the presence of a thin lamellar phase evenly distributed throughout the structure. matrix phase, and there is no eutectic phase. The element distribution and EMF point analysis of the lamellar phase revealed a relatively high content of Cr (yellow) and C (green) in this phase, which again indicates the decomposition of SiC during melting and the interaction of the released carbon with the chromium effect. The VEA matrix forms a lamellar carbide phase. The distribution of elements and point analysis of the matrix phase showed that most of the cobalt, iron, nickel and silicon are present in the matrix phase.
(a) SEM image of sample S-2, (b) enlarged image, (c) element map, (d) EMF results at indicated locations.
SEM studies of C-3 composites revealed the presence of new phases in addition to the carbide and matrix phases. The elemental map (Fig. 4c) and EMF point analysis (Fig. 4d) show that the new phase is rich in nickel, cobalt, and silicon.
(a) SEM image of sample S-3, (b) enlarged image, (c) element map, (d) EMF results at indicated locations.
The results of the SEM and EMF analysis of the C-4 composite are shown in Figs. 5. In addition to the three phases observed in composite C-3, the presence of graphite nodules was also found. The volume fraction of the silicon-rich phase is also higher than that of the C-3 composite.
(a) SEM image of sample S-4, (b) enlarged image, (c) element map, (d) EMF results at indicated locations.
The results of the SEM and EMF spectra of composites S-5 and S-6 are shown in Figures 1 and 2. 6 and 7, respectively. In addition to a small number of spheres, the presence of graphite flakes was also observed. Both the number of graphite flakes and the volume fraction of the silicon-containing phase in the C-6 composite are greater than in the C-5 composite.
(a) SEM image of sample C-5, (b) enlarged view, (c) elemental map, (d) EMF results at indicated locations.
(a) SEM image of sample S-6, (b) enlarged image, (c) element map, (d) EMF results at indicated locations.
Crystal structure characterization of HEA composites was also performed using XRD measurements. The result is shown in Figure 8. In the diffraction pattern of the base WEA (S-0), only the peaks corresponding to the fcc phase are visible. X-ray diffraction patterns of composites C-1, C-2, and C-3 revealed the presence of additional peaks corresponding to chromium carbide (Cr7C3), and their intensity was lower for samples C-3 and C-4, which indicated that also with the data EMF for these samples. Peaks corresponding to Co/Ni silicides were observed for samples S-3 and S-4, again consistent with the EDS mapping results shown in Figures 2 and 3. As shown in Figure 3 and Figure 4. 5 and S-6 peaks were observed corresponding to graphite.
Both microstructural and crystallographic characteristics of the developed composites indicated decomposition of the added SiC. This is due to the presence of chromium in the VEA matrix. Chromium has a very strong affinity for carbon 54.55 and reacts with free carbon to form carbides, as indicated by the observed decrease in the chromium content of the matrix. Si passes into the fcc phase due to the dissociation of SiC56. Thus, an increase in the addition of SiC to the base HEA led to an increase in the amount of the carbide phase and the amount of free Si in the microstructure. It has been found that this additional Si is deposited in the matrix at low concentrations (in composites S-1 and S-2), while at higher concentrations (composites S-3 to S-6) it results in additional cobalt deposition/. nickel silicide. The standard enthalpy of formation of Co and Ni silicides, obtained by direct synthesis high-temperature calorimetry, is -37.9 ± 2.0, -49.3 ± 1.3, -34.9 ± 1.1 kJ mol -1 for Co2Si, CoSi and CoSi2, respectively, while these values are – 50.6 ± 1.7 and – 45.1 ± 1.4 kJ mol-157 for Ni2Si and Ni5Si2, respectively. These values are lower than the heat of formation of SiC, indicating that the dissociation of SiC leading to the formation of Co/Ni silicides is energetically favorable. In both S-5 and S-6 composites, additional free silicon was present, which was absorbed beyond the formation of silicide. This free silicon has been found to contribute to the graphitization observed in conventional steels58.
The mechanical properties of the developed ceramic-reinforced composites based on HEA are investigated by compression tests and hardness tests. The stress-strain curves of the developed composites are shown in Figs. 9a, and in Fig. 9b shows a scatterplot between specific yield strength, yield strength, hardness, and elongation of the developed composites.
(a) Compressive strain curves and (b) scatterplots showing specific yield stress, yield strength, hardness and elongation. Note that only specimens S-0 to S-4 are shown, as specimens S-5 and S-6 contain significant casting defects.
As seen in fig. 9, the yield strength increased from 136 MPa for the base VES (C-0) to 2522 MPa for the C-4 composite. Compared to the basic WPP, the S-2 composite showed a very good elongation to failure of about 37%, and also showed significantly higher yield strength values (1200 MPa). The excellent combination of strength and ductility of this composite is due to the improvement in the overall microstructure, including the uniform distribution of fine carbide lamellae throughout the microstructure, which is expected to inhibit dislocation movement. The yield strengths of C-3 and C-4 composites are 1925 MPa and 2522 MPa, respectively. These high yield strengths can be explained by the high volume fraction of cemented carbide and silicide phases. However, the presence of these phases also resulted in an elongation at break of only 7%. The stress-strain curves of the base composites CoCrFeNi HEA (S-0) and S-1 are convex, indicating activation of the twinning effect or TRIP59,60. Compared to sample S-1, the stress-strain curve of sample S-2 has a concave shape at a strain of about 10.20%, which means that the normal dislocation slip is the main deformation mode of the sample in this deformed state60,61. However, the hardening rate in this specimen remains high over a large strain range, and at higher strains a transition to convexity is also visible (although it cannot be ruled out that this is due to the failure of lubricated compressive loads). ). Composites C-3 and C-4 have only limited plasticity due to the presence of higher volume fractions of carbides and silicides in the microstructure. Compression tests of samples of composites C-5 and C-6 were not carried out due to significant casting defects on these samples of composites (see Fig. 10).
Stereomicrographs of casting defects (indicated by red arrows) in samples of composites C-5 and C-6.
The results of measuring the hardness of VEA composites are shown in Figs. 9b. The base WEA has a hardness of 130±5 HV, and samples S-1, S-2, S-3 and S-4 have hardness values of 250±10 HV, 275±10 HV, 570±20 HV and 755±20 HV. The increase in hardness was in good agreement with the change in yield strength obtained from compression tests and was associated with an increase in the amount of solids in the composite. The calculated specific yield strength based on the target composition of each sample is also shown in fig. 9b. In general, the best combination of yield strength (1200 MPa), hardness (275 ± 10 HV), and relative elongation to failure (~37%) is observed for composite C-2.
Comparison of the yield strength and relative elongation of the developed composite with materials of different classes is shown in Fig. 11a. Composites based on CoCrFeNi in this study showed high elongation at any given stress level62. It can also be seen that the properties of the HEA composites developed in this study lie in the previously unoccupied region of the plot of yield strength versus elongation. In addition, the developed composites have a wide range of combinations of strength (277 MPa, 1200 MPa, 1925 MPa and 2522 MPa) and elongation (>60%, 37%, 7.3% and 6.19%). Yield strength is also an important factor in the selection of materials for advanced engineering applications63,64. In this regard, the HEA composites of the present invention exhibit an excellent combination of yield strength and elongation. This is because the addition of low density SiC results in composites with high specific yield strength. The specific yield strength and elongation of HEA composites are in the same range as HEA FCC and refractory HEA, as shown in Fig. 11b. The hardness and yield strength of the developed composites are in the same range as for massive metallic glasses65 (Fig. 11c). Massive metallic glasses (BMS) are characterized by high hardness and yield strength, but their elongation is limited66,67. However, the hardness and yield strength of some of the HEA composites developed in this study also showed significant elongation. Thus, it was concluded that the composites developed by VEA have a unique and sought-after combination of mechanical properties for various structural applications. This unique combination of mechanical properties can be explained by the uniform dispersion of hard carbides formed in situ in the FCC HEA matrix. However, as part of the goal of achieving a better combination of strength, microstructural changes resulting from the addition of ceramic phases must be carefully studied and controlled to avoid casting defects, such as those found in S-5 and S-6 composites, and ductility. gender.
The results of this study were compared with various structural materials and HEAs: (a) elongation versus yield strength62, (b) specific yield stress versus ductility63 and (c) yield strength versus hardness65.
The microstructure and mechanical properties of a series of HEA-ceramic composites based on the HEA CoCrFeNi system with the addition of SiC have been studied and the following conclusions have been drawn:
High entropy alloy composites can be successfully developed by adding SiC to CoCrFeNi HEA using the arc melting method.
SiC decomposes during arc melting, leading to the formation in situ of carbide, silicide and graphite phases, the presence and volume fraction of which depend on the amount of SiC added to the base HEA.
HEA composites exhibit many excellent mechanical properties, with properties that fall into previously unoccupied areas on the yield strength versus elongation plot. The yield strength of the HEA composite made using 6 wt% SiC was more than eight times that of base HEA while maintaining 37% ductility.
The hardness and yield strength of HEA composites are in the range of bulk metallic glasses (BMG).
The findings suggest that high-entropy alloy composites represent a promising approach to achieving an excellent combination of metal-mechanical properties for advanced structural applications.
Post time: Jul-12-2023