Effect of High-Pressure Torsion on the Microstructure and Magnetic Properties of Nanocrystalline CoCrFeNiGax (x = 0.5, 1.0) High Entropy Alloys

In our search for an optimum soft magnet with excellent mechanical properties which can be used in applications centered around “electro mobility”, nanocrystalline CoCrFeNiGax (x = 0.5, 1.0) bulk high entropy alloys (HEA) were successfully produced by spark plasma sintering (SPS) at 1073 K of HEA powders produced by high energy ball milling (HEBM). SPS of non-equiatomic CoCrFeNiGa0.5 particles results in the formation of a single-phase fcc bulk HEA, while for the equiatomic CoCrFeNiGa composition a mixture of bcc and fcc phases was found. For both compositions SEM/EDX analysis showed a predominant uniform distribution of the elements with only a small number of Cr-rich precipitates. High pressure torsion (HPT) of the bulk samples led to an increased homogeneity and a grain refinement: i.e., the crystallite size of the single fcc phase of CoCrFeNiGa0.5 decreased by a factor of 3; the crystallite size of the bcc and fcc phases of CoCrFeNiGa—by a factor of 4 and 10, respectively. The lattice strains substantially increased by nearly the same extent. After HPT the saturation magnetization (Ms) of the fcc phase of CoCrFeNiGa0.5 and its Curie temperature increased by 17% (up to 35 Am2/kg) and 31.5% (from 95 K to 125 K), respectively, whereas the coercivity decreased by a factor of 6. The overall Ms of the equiatomic CoCrFeNiGa decreased by 34% and 55% at 10 K and 300 K, respectively. At the same time the coercivity of CoCrFeNiGa increased by 50%. The HPT treatment of SPS-consolidated HEAs increased the Vickers hardness (Hv) by a factor of two (up to 5.632 ± 0.188) only for the non-equiatomic CoCrFeNiGa0.5, while for the equiatomic composition, the Hv remained unchanged (6.343–6.425 GPa).


Introduction
In general, metallic materials currently available for most applications are based on one or two principal elements: Cu-based alloys, Ni-based superalloys, Fe-based steels, etc. [1,2]. Developing new alloys has mostly been focused on microstructure/surface modification and adding alloying elements to improve specific properties.
In 2004 Yeh et al. [3] showed that multi principal component alloys containing 5 or more principal elements in equiatomic or nearly equiatomic amounts (ranging between 5 and 35 at. %) could be processed to form simple solid solutions under appropriate conditions, called high entropy alloys (HEAs). Their high configurational entropy of mixing (especially at higher temperatures) improves the stability of chemically disordered
HEBM was performed in a water-cooled planetary ball mill Activator-2S using stainless steel cylindrical jars and steel balls (7 mm in diameter). In all cases, the ball/powder weight ratio was 20:1. The vial was evacuated and then filled with Ar gas up to 4 bars. The HEBM was run at a rotating speed of the sun wheel and the grinding drums at 900 and 1800 rpm, respectively. The milling time (t) in Ar reached 180 min. An additional t = 10 min (in C 3 H 7 OH) of milling after HEBM in Ar was applied.
The non-milled and milled CoCrFeNiGa x (x = 0.5, 1.0) powders were SPS-consolidated in vacuum in a Labox 650 facility (Sinter Land, Japan). The powder mixture was placed into a cylindrical graphite die (inner diameter 12.7 mm) and uniaxially compressed at 10-50 MPa. The sample was heated at a rate of 100 K/min up to 1073 K by passing rectangular pulses of electric current through the sample. The dwell time at sintering temperature was 10 min. SPS-produced disks were 2-3 mm thick and 12.7 mm in diameter.
A set of SPS-consolidated samples (disks of 10 mm in diameter) were severely deformed by 10 rotations (6 rotations for the HEBM sample CoCrFeNiGa 0.5 , Table 1) in high pressure torsion at room temperature (RT). The process was conducted at a rate of 1 rpm and an applied contact pressure of 5.1 GPa, using water cooling to keep the temperature close to RT (30-40 • C). A schematic diagram of the processes is presented in Figure 1.

Materials and Methods
CoCrFeNiGax (x = 0.5, 1.0) HEA powders were prepared by high energy ball milling (HEBM) of commercial powders of Co (99.97% pure, particle size ~45 µm), Cr (99.7% pure, particle size 10-30 µm), Fe (99.96% pure, particle size < 150 µm), Ni (99.5% pure, particle size 45-60 µm), and Ga (99.99% pure, ingot) taken in aliquot amounts. For the equiatomic CoCrFeNiGa composition, the amount of 20 at. % of each element was taken, whereas for the CoCrFeNiGa0.5 alloy-as follows (at. HEBM was performed in a water-cooled planetary ball mill Activator-2S using stainless steel cylindrical jars and steel balls (7 mm in diameter). In all cases, the ball/powder weight ratio was 20:1. The vial was evacuated and then filled with Ar gas up to 4 bars. The HEBM was run at a rotating speed of the sun wheel and the grinding drums at 900 and 1800 rpm, respectively. The milling time (t) in Ar reached 180 min. An additional t = 10 min (in C3H7OH) of milling after HEBM in Ar was applied.
The non-milled and milled CoCrFeNiGax (x = 0.5, 1.0) powders were SPS-consolidated in vacuum in a Labox 650 facility (Sinter Land, Japan). The powder mixture was placed into a cylindrical graphite die (inner diameter 12.7 mm) and uniaxially compressed at 10-50 MPa. The sample was heated at a rate of 100 K/min up to 1073 K by passing rectangular pulses of electric current through the sample. The dwell time at sintering temperature was 10 min. SPS-produced disks were 2-3 mm thick and 12.7 mm in diameter.
A set of SPS-consolidated samples (disks of 10 mm in diameter) were severely deformed to 10 rotations (6 rotations for the HEBM sample CoCrFeNiGa0.5, Table 1) by high pressure torsion at room temperature (RT). The process was conducted at a rate of 1 rpm and an applied contact pressure of 5.1 GPa, using water cooling to keep the temperature close to RT (30-40 °C). A schematic diagram of the processes is presented in Figure 1. The preparation conditions for both compositions are summarized in Table 1.  The preparation conditions for both compositions are summarized in Table 1. Initial and milled powders, SPS-consolidated and HPT-deformed samples were characterized by X-ray diffraction (XRD) using Fe-Kα radiation to distinguish structures of Co, Fe, and Ni (DRON-3M diffractometer, Fe-Kα radiation with λ = 0.19374 nm, 2θ = 40-120 • ). Crystalline phases were identified using Crystallographica Search-Match 2.1 software and ICDD PDF2 database. To determine crystal cell parameters, crystallite size and strains we used the PDWin 6 software (NPP Bourevestnik, Saint-Petersburg, Russia, 2010) [31]. For the calculations of the crystalline phase content, the Rietveld method was used.
SPS-consolidated and HPT-processed samples were cross sectioned along the diameter and polished to analyze their microstructure and chemical composition by scanning electron microscopy (SEM, with operating voltage 15 kV) with operating voltage 15 kV along with energy dispersive spectroscopy (EDX) (Zeiss Ultra+ microscope + Oxford Inca spectrometer, Aztec 2.1 software). Magnetic properties of the samples were determined using a Quantum Design DynaCool Physical Property Measurement system (PPMS) at temperatures in the range of 5-390 K under external magnetic fields up to 9 Tesla.
The microhardness of SPS-consolidated/HPT-deformed CoCrFeNiGa x (x = 0.5, 1.0) samples was measured using Vickers hardness tests with an Emco-Test DuraScan 70 (Austria) under an applied load of 4.9 N (HV 0.5 ). We focused our study on tuning the structure and magnetic properties by HPT for bulk CoCrFeNiGa x (x = 0.5, 1.0) HEAs produced by SPS from HEBM powders (190 min at 900/1800 rpm), as bulk HEAs from elemental powders Co, Cr, Fe, Ni, and Ga ingots by SPS could not be fabricated (for details, see Section S1 in the Supplementary File).

Results and Discussion
For SPS consolidation and subsequent HPT treatment the CoCrFeNiGa 0.5 and CoCr-FeNiGa HEA powders with a fcc lattice parameters (calculated as a = 3.610 ± 0.005 Å and a = 3.632 ± 0.007 Å, respectively), and a uniform distribution of principal elements obtained by HEBM were used [19]. The details of their structural and magnetic behavior after HEBM are described in [19]. Figure 2 shows XRD data of SPS-consolidated (dash red) and HPT-deformed (blue) CoCrFeNiGa 0.5 HEAs produced from HEA powders (black).
SPS-consolidation of HEBM-processed HEA powder ( Figure 2, black) at 1073 K leads to an increase in crystallinity of the fcc phase ( Figure 2, dash red). The XRD (111), (200), and (220) peaks become narrower after sintering due to an increase in the crystallite size and decrease in lattice strains caused by HEBM. Other additional phases were not observed. The subsequent HPT treatment of the SPS-consolidated HEA ( Figure 2, dash red) led to a decrease in intensity of XRD (111), (200), and (220) peaks (Figure 2, blue) due a grain refinement and a substantial increase in lattice strains of the CoCrFeNiGa 0.5 fcc singlephase ( Table 2). The phase transformation has not been occurred in HPT-deformed HEA (Figure 2, blue).
The SEM/EDX results for both the SPS-consolidated and the subsequently HPTdeformed CoCrFeNiGa 0.5 HEAs show that the elements of Co, Cr, Fe, Ni, and Ga are primarily uniformly distributed on a micro-scale (Figure 3a,b). Black Cr-rich precipitates were also detected by EDX. Some traces of laminar flow structures and increased alloy homogeneity were observed caused by HPT (Figure 3b). SPS consolidation of the equiatomic CoCrFeNiGa HEA powder (Figure 4, black) at 1073 K led to partial decomposition of the fcc structure. XRD data showed that SPSconsolidated material (Figure 4, dash red) contained a mixture of bcc and fcc phases, those crystallite sizes decreased, and lattice strains increased after HPT treatment ( Figure 4, blue, Table 2). It is important to note that there is no evidence for appearance of a new phase and the occurrence of the phase transformation during HPT processing.  The SEM/EDX results for both the SPS-consolidated and the subsequently HPT-deformed CoCrFeNiGa0.5 HEAs show that the elements of Co, Cr, Fe, Ni, and Ga are primarily uniformly distributed on a micro-scale (Figure 3a,b). Black Cr-rich precipitates were also detected by EDX. Some traces of laminar flow structures and increased alloy homogeneity were observed caused by HPT (Figure 3b).  The SEM/EDX results of SPS-consolidated and HPT-deformed CoCrFeNiGa HEAs ( Figure 5) show a uniform equiatomic distribution of the principal elements. Almost spherical grains (Figure 5a) transform into elongated grains ( Figure 5b). The presence of few Cr-rich precipitates for both HEAs was also detected; however, it is evident that the interfaces of these precipitates are less sharp after the HPT process.
Our previous work [19] showed that the crystallite sizes of single-phase fcc HEBMprocessed HEA CoCrFeNiGa x (x = 0.5, 1.0) powders were around 9 nm. The microstrains were 0.12% and 0.36% for CoCrFeNiGa and CoCrFeNiGa 0.5 , respectively.  SPS consolidation of the equiatomic CoCrFeNiGa HEA powder ( Figure 4, black) at 1073 K led to partial decomposition of the fcc structure. XRD data showed that SPS-consolidated material ( Figure 4, dash red) contained a mixture of bcc and fcc phases, those crystallite sizes decreased, and lattice strains increased after HPT treatment ( Figure 4, blue, Table 2). It is important to note that there is no evidence for creating a new phase and the occurrence of the phase transformation during HPT processing.   (Figure 5b). The presence of few Cr-rich precipitates for both HEAs was also detected; however, it is evident that the interfaces of these precipitates are less sharp after the HPT process.  The crystallite size of the single-phase fcc CoCrFeNiGa0.5 HEA powder [19] is seen to grow by a factor of 6 after SPS-consolidation and decreases by a factor of 3 after HPT treatment which is typical for such processes. The lattice strains show opposite behaviorreduced after SPS-consolidation at 1073 K and increased again during HPT deformation ( Table 2).
For the equiatomic SPS-consolidated CoCrFeNiGa HEA the crystallite size of both bcc and fcc phases decreases by a factor of 4 and 10 after HPT, respectively. Microstrains of fcc and bcc phases in HPT-deformed CoCrFeNiGa HEA increased to nearly the same extent (see Table 2). It has to be mentioned that the measured strain of 1.73% after HPT is probably not the true strain but the maximum strain due to defects and additional broadening mechanisms. A slight increase in crystal lattice for both fcc and bcc structures in the The crystallite size and strain in bulk CoCrFeNiGa x (x = 0.5, 1.0) was derived from the line width analysis of XRD peaks. The pseudo-Voigt function was used for fitting of the XRD peak profile, and a Si standard was used to correct for instrumental broadening. Crystallite size, lattice strain, and lattice parameters for bcc and fcc phases were calculated for SPS-consolidated and HPT-deformed CoCrFeNiGa x (x = 0.5, 1.0) HEAs using the method of second central moments [31]. The results of the calculations are summarized in Table 2.
The crystallite size of the single-phase fcc CoCrFeNiGa 0.5 HEA powder [19] increases by a factor of 6 after SPS-consolidation and decreases by a factor of 3 after HPT treatment which is typical for such processes. The lattice strains show opposite behavior-reduced after SPS-consolidation at 1073 K and increased again during HPT deformation (Table 2).
For the equiatomic SPS-consolidated CoCrFeNiGa HEA the crystallite size of both bcc and fcc phases decreases by a factor of 4 and 10 after HPT, respectively. Microstrains of fcc and bcc phases in HPT-deformed CoCrFeNiGa HEA increased to nearly the same extent (see Table 2). It has to be mentioned that the measured strain of 1.73% after HPT is probably not the true strain but the maximum strain due to defects and additional broadening mechanisms. A slight increase in crystal lattice for both fcc and bcc structures in the bulk HEAs (Table 2) after HPT can be attributed to the partial dissolution of the precipitates previously observed in SPS-consolidated HEAs (Figures 3 and 5).
The values of volume-to-volume ratio of bcc (~84%, R wp = 6.87%) to fcc (~16%, R wp = 6.87%) solid solutions in SPS-consolidated CoCrFeNiGa alloy slightly changed after the HPT deformation: the amount of a bcc phase decreased (~75%, R wp = 7.76%), while the amount of fcc one increased (~25%, R wp = 7.76). This phase transformation can be linked to the deformation during the HPT process, and was recently also reported for HEAs in the literature [32][33][34].
SEM-BSE micrographs ( Figure 6) at higher magnification (×40,000) show that SPSconsolidated CoCrFeNiGa 0.5 HEA (Figure 6a) undergoes nanostructuring (drastic decrease in grain size) and an increased homogeneity upon HPT without any significant severe grain  (Figure 6b). For equiatomic CoCrFeNiGa bulk (Figure 6c), a lamellar microstructure with a pronounced two-phase separation was observed after HPT. The SEM-BSE results agree well with the XRD data (Figures 3 and 4).
SEM-BSE micrographs ( Figure 6) at higher magnification (×40,000) show that SPSconsolidated CoCrFeNiGa0.5 HEA (Figure 6a) undergoes nanostructuring (drastic decrease in grain size) and an increased homogeneity upon HPT without any significant severe grain flow (Figure 6b). For equiatomic CoCrFeNiGa bulk (Figure 6c), a lamellar microstructure with a pronounced two-phase separation was observed after HPT. The SEM-BSE results agree well with the XRD data (Figures 3 and 4).
In Figure 6 the Ga-rich areas (Z = 31) look brighter than the other elements: Cr (Z = 24), Co (Z = 27), Fe (Z = 26), and Ni (Z = 28). We assume that the darker areas on Figure 6d correspond to the Ga-depleted phase with the fcc structure ( Figure 4) and the brighter regions to the Ga-rich bcc phase. Figure 6. SEM-BSE images of (a,b)-bulk CoCrFeNiGa0.5 HEA before and after HPT, respectively; (c,d)-bulk CoCrFeNiGa HEA before and after HPT, respectively.

CoCrFeNiGa0.5 HEAs
HPT deformation of the SPS-consolidated CoCrFeNiGa0.5 HEA leads to an increase in the saturation magnetization and the Curie temperature (Figure 7). A very smooth transition from a ferro-to a paramagnetic state is observed. The magnetization of both samples (Figure 7a) decreases already at temperatures much lower than the approximate Curie temperature Tc. This behavior of M(T) differs from classical ferromagnets where a significant decrease in the magnetization starts at T > 2/3 Tc followed by a rapid loss of spontaneous magnetization at the Curie temperature. Smoothing of the transition across such a wide temperature range indicates magnetic inhomogeneities which is possibly due to a Figure 6. SEM-BSE images of (a,b)-bulk CoCrFeNiGa 0.5 HEA before and after HPT, respectively; (c,d)-bulk CoCrFeNiGa HEA before and after HPT, respectively.
In Figure 6 the Ga-rich areas (Z = 31) look brighter than the other elements: Cr (Z = 24), Co (Z = 27), Fe (Z = 26), and Ni (Z = 28). We assume that the darker areas on Figure 6d correspond to the Ga-depleted phase with the fcc structure ( Figure 4) and the brighter regions to the Ga-rich bcc phase.

CoCrFeNiGa 0.5 HEAs
HPT deformation of the SPS-consolidated CoCrFeNiGa 0.5 HEA leads to an increase of the saturation magnetization and the Curie temperature (Figure 7). A very broad transition from a ferro-to a paramagnetic state is observed. The magnetization of both samples (Figure 7a) decreases already at temperatures much lower than the approximate Curie temperature T c . This behavior of M(T) differs from classical ferromagnets where a significant decrease in the magnetization starts at T > 2/3 T c followed by a rapid loss of spontaneous magnetization at the Curie temperature. Smoothing of the transition across such a wide temperature range indicates magnetic inhomogeneities which are possibly due to a broad distribution of local magnitudes and signs of the exchange interactions as a consequence of disorder on the atomic scale.
An apparent Curie temperature can be determined according to the Landau theory as a minimum of the first derivative dM/dT (Figure 7b). This point marks a crossover from an exchange-ordered to a field-ordered magnetic state [35]. Using this model, we find that HPT results in an increase of this apparent Curie temperature from 95 K to 125 K.  An apparent Curie temperature can be determined according to the Landau theory as a minimum of the first derivative dM/dT (Figure 7b). This point marks a crossover from an exchange-ordered to a field-ordered magnetic state [35]. Using this model, we find that HPT results in an increase in this apparent Curie temperature from 95 K to 125 K.
On the other hand, a "paramagnetic ordering temperature" can be estimated from the Curie-Weiss law as the temperature at which the paramagnetic inverse susceptibility 1/χ is zero.
Well above in the paramagnetic state, the temperature dependence of the susceptibility is described by the Curie-Weiss law: = , where C is the Curie constant = µ µ . Here n is the volume density of atoms, μ0 is the permeability of free space, is Boltzmann's constant and µ is the squared effective magnetic moment per atom.
An effective magnetic moment per atom can be defined using the formula: = .
A fit according to the Curie-Weiss law (green dash lines in Figure 7c) yields the average moment per atom 1.52µ for both samples: before and after HPT treatment and a "paramagnetic" Curie temperature, , of 125 K and 185 K, respectively. The atomic density ~8.57 × 10 m was estimated based on the results of XRD studies. Note that the Curie temperature, , is much higher than the temperature we have determined earlier ( Figure  7b). We assume that above "Tc" the long-range ferromagnetic correlations observed as M(T) are stabilized by an external magnetic field. Conventionally, the average magnetic moments per atom can be determined from the low temperature saturation magnetization. We find ~0.31µ and ~0.36µ for SPS consolidated alloys before and after HPT, respectively. The large discrepancy between the magnetic moments determined from the low temperature magnetization (0.31µ ) and the Curie-Weiss-law (1.52µ ) indicates the presence of antiferromagnetic exchange interactions which are thermally overcome in the paramagnetic regime. Thus, we may conclude that Cr is antiferromagnetically correlated to neighboring elements in the equatomic CrFe-CoNi alloy [36,37]. Note that reflecting the dominance of thermal energy over the exchange correlations is much higher than Tc. Also, a "paramagnetic ordering temperature" θ C can be estimated from the Curie-Weiss law as the temperature at which the paramagnetic inverse susceptibility 1/χ is zero.
Well above θ C in the paramagnetic state, the temperature dependence of the susceptibility is described by the Curie-Weiss law: Here n is the volume density of atoms, µ 0 is the permeability of free space, k B is Boltzmann's constant and µ 2 e f f is the squared effective magnetic moment per atom. An effective magnetic moment per atom can be defined using the formula: µ e f f = nCµ 0 3k B . A fit according to the Curie-Weiss law (green dash lines in Figure 7c) yields the average moment per atom 1.52µ B for both samples: before and after HPT treatment and a "paramagnetic" Curie temperature, θ C , of 125 K and 185 K, respectively. The atomic density ∼ 8.57 × 10 28 m −3 was estimated based on the results of XRD studies. Note that the "paramagnetic" Curie temperature, θ C , is much higher than the "ferromagnetic" temperature we have determined earlier (Figure 7b). We assume that above "T c " the long-range ferromagnetic correlations observed as M(T) are stabilized by an external magnetic field.
Conventionally, the average magnetic moments per atom can be determined from the low temperature saturation magnetization. We find~0.31µ B and~0.36µ B for SPS consolidated alloys before and after HPT, respectively. The large discrepancy between the magnetic moments determined from the low temperature magnetization (0.31µ B ) and the Curie-Weiss-law (1.52µ B ) indicates the presence of antiferromagnetic exchange interactions which are thermally overcome in the paramagnetic regime. Thus, we may conclude that Cr is antiferromagnetically correlated to neighboring elements in the equiatomic CrFeCoNi alloy [36,37]. Note that θ C reflecting the dominance of thermal energy over the exchange correlations is much higher than T c .
How to interpret the discrepancy between the magnetic moments determined from the saturation magnetization and in the paramagnetic regime? In the ferromagnetic state we determine the average magnetic moment given by the band structure, whereas in the paramagnetic mode we assume the so-called local magnetic moment, which may in turn deviate from the present atomic moment due to ferromagnetic interatomic correlations still existing above the Curie temperature [38].
In addition, antiferromagnetic exchange interatomic interactions are expected in CoCr-FeNiGa alloy, predominantly at the Cr sites, similar to those observed in the equiatomic CrFeCoNi alloy [36,37].
The deviation of 1/χ(T) above 250 K from the paramagnetic behavior for the asprepared CoCrFeNiGa 0.5 HEA can be attributed to the presence of superparamagnetic precipitates with a blocking temperature of 250 K which give rise to an effective paramagnetic susceptibility. HPT of the CoCrFeNiGa 0.5 results in the dissolution of these precipitates and the formation of a more homogeneous alloy as was observed in SEM/EDX studies (Figure 3b). The field dependence of the magnetization of SPS-consolidated and HPT-deformed CoCrFeNiGa 0.5 HEAs (Figure 8b) confirms this interpretation: dissolution of the precipitates which act as pinning centers for domain walls results in magnetic softening of the alloy after HPT processing.
we determine the average magnetic moment given by the band structure, whereas in the paramagnetic mode we assume the so-called local magnetic moment, which may in turn deviate from the present atomic moment due to ferromagnetic interatomic correlations still existing above the Curie temperature [38].
In addition, antiferromagnetic exchange interatomic interactions are expected in CoCrFeNiGa alloy, predominantly at the Cr sites, similar to those observed in the equiatomic CrFeCoNi alloy [36,37].
The other point to discuss: the deviation of 1/χ(T) above 250 K from the paramagnetic behavior for the as-prepared CoCrFeNiGa0.5 HEA can be attributed to the presence of superparamagnetic precipitates with a blocking temperature of 250 K which give rise to an effective paramagnetic susceptibility. HPT of the CoCrFeNiGa0.5 results in the dissolution of these precipitates and the formation of a more homogeneous alloy as was observed in SEM/EDX studies (Figure 3b). The field dependence of the magnetization of SPS-consolidated and HPT-deformed CoCrFeNiGa0.5 HEAs (Figure 8b) confirms this interpretation: dissolution of the precipitates which act as pinning centers for domain walls results in magnetic softening of the alloy after HPT processing.  Figure 9 shows temperature dependences of magnetization for SPS-consolidated and HPT-deformed equiatomic CoCrFeNiGa HEAs. In contrast to CoCrFeNiGa0.5 the equiatomic alloy shows a higher saturation magnetization and a Tc well above 390 K. The increase in M(T) below 40 K approximately indicates the presence of an additional ferromagnetic phase with an ordering temperature of ~22 K and a magnetization which is about 2 Am 2 /kg, i.e., ~4.5% of the other phase (blue curve).

CoCrFeNiGa High Entropy Alloy
HPT processing results in a significant decrease in the overall magnetization. The low-temperature ferromagnetic phase manifests itself more clearly, and its Curie temperature increases up to 40 K. A larger magnetic inhomogeneity of the high-Curie temperature phase is evidenced by the smoother transition to the paramagnetic state.  Figure 9 shows temperature dependences of magnetization for SPS-consolidated and HPT-deformed equiatomic CoCrFeNiGa HEAs. In contrast to CoCrFeNiGa 0.5 the equiatomic alloy shows a higher saturation magnetization and a T c well above 390 K. The increase in M(T) below 40 K approximately indicates the presence of an additional ferromagnetic phase with an ordering temperature of~22 K and a magnetization which is about 2 Am 2 /kg, i.e.,~4.5% of the other phase (blue curve).

CoCrFeNiGa High Entropy Alloy
HPT processing results in a significant decrease in the overall magnetization. The lowtemperature ferromagnetic phase manifests itself more clearly, and its Curie temperature increases up to 40 K. A larger magnetic inhomogeneity of the high-Curie temperature phase is evidenced by the smoother transition to the paramagnetic state.
HPT processing of the equiatomic CoCrFeNiGa HEA results in a decrease in the saturation magnetization by 34% and 55% (see Table 3) at 10 K and 300 K, respectively ( Figure 10). The decrease in M s is due to the increased volume fraction of the Ga-depleted fcc phase, which has a lower magnetization (Figure 8a) compared to the bcc Ga-enriched phase [17].
Note that the coercivity H c of both samples is higher at room temperature than at 10 K. Above the Curie temperature of the fcc phase, ferromagnetic exchange between bcc nanocrystalline ferromagnetic grains is weakened resulting in magnetic hardening [39]. H c (300 K) for the HPT deformed alloy is larger by 50% in comparison to the SPS-consolidated one. This increase can be attributed to the introduction of defects which act as pinning centers for domain wall movement. HPT processing of the equiatomic CoCrFeNiGa HEA results in a decrease in the saturation magnetization by 34% and 55% (see Table 3) at 10 K and 300 K, respectively (Figure 10). The decrease in Ms is due to the increased volume fraction of the Ga-depleted fcc phase, which has a lower magnetization (Figure 8a) compared to the bcc Ga-enriched phase [17].     HPT processing of the equiatomic CoCrFeNiGa HEA results in a decrease in the saturation magnetization by 34% and 55% (see Table 3) at 10 K and 300 K, respectively (Figure 10). The decrease in Ms is due to the increased volume fraction of the Ga-depleted fcc phase, which has a lower magnetization (Figure 8a) compared to the bcc Ga-enriched phase [17].

Mechanical Properties (Vickers Microhardness)
The Vickers hardness (H v ) for the SPS-consolidated and HPT-deformed CoCrFeNiGa x (x = 0.5, 1.0) samples from elemental powder blends showed the lowest H v with significant variations for both compositions (for details, see Section S2 in the Supplementary File).
The dependence of the Vickers hardness (H v ) before and after HPT for SPS-consolidated samples obtained from HEBM CoCrFeNiGa x (x = 0.5, 1.0) powders [19] are presented in Figure 11. nanocrystalline ferromagnetic grains is weakened resulting in magnetic hardening [39]. Hc (300 K) for the HPT deformed alloy is larger by 50% in comparison to the SPS-consolidated one. This increase can be attributed to the introduction of defects which act as pinning centers for domain wall movement.

Mechanical Properties (Vickers Microhardness)
The Vickers hardness (Hv) for the SPS-consolidated and HPT-deformed CoCrFeNi-Gax (x = 0.5, 1.0) samples from elemental powder blends showed the lowest Hv with significant variations for both compositions (for details, see Section S2 in the Supplementary File).
The dependence of the Vickers hardness (Hv) before and after HPT for SPS-consolidated samples obtained from HEBM CoCrFeNiGax (x = 0.5, 1.0) powders [19] are presented in Figure 11.  Figure 11. Dependence of the Vickers hardness on the radial distance from the center for SPS-consolidated (red), and subsequently HPT-deformed (dark blue) CoCrFeNiGax (x = 0.5, 1.0) samples from HEBM powder mixtures.
The Hv of the SPS-consolidated CoCrFeNiGa and CoCrFeNiGa0.5 HEAs (HEBM processed) is 3 and 5 times higher than for bulk samples sintered from elemental powder blends (see Section S2 in the Supplementary File), respectively. It is worth noting, that the subsequent HPT treatment of SPS-consolidated HEAs increased Hv by a factor of two (up to 5.632 ± 0.188 GPa) only for the non-equiatomic CoCrFeNiGa0.5, while for the equiatomic composition, Hv remained unchanged. The high hardness for SPS-consolidated equiatomic CoCrFeNiGa HEAs may be attributed to the mixture of fcc and bcc phases, and Cr-rich precipitates which hinders the gliding of dislocations at grain boundaries. Subsequent refinement of the microstructure by applied shear strain (HPT) does not significantly change the microstructure of grains. Only a slight elongation of the grains was observed ( Figure 11). The H v of the SPS-consolidated CoCrFeNiGa and CoCrFeNiGa 0.5 HEAs (HEBM processed) is 3 and 5 times higher than for bulk samples sintered from elemental powder blends (see Section S2 in the Supplementary File), respectively. It is worth noting, that the subsequent HPT treatment of SPS-consolidated HEAs increased H v by a factor of two (up to 5.632 ± 0.188 GPa) only for the non-equiatomic CoCrFeNiGa 0.5 , while for the equiatomic composition, H v remained unchanged. The high hardness for SPS-consolidated equiatomic CoCrFeNiGa HEAs may be attributed to the mixture of fcc and bcc phases, and Cr-rich precipitates which hinders the gliding of dislocations and lattice planes at grain boundaries. Subsequent refinement of the microstructure by applied shear strain (HPT) does not significantly change the microstructure of grains. Only a slight elongation of the grains was observed ( Figure 11).
As expected, an insignificant gradient in hardness along the radius of the disk is visible as for SPS-consolidated as for HPT-deformed samples ( Figure 11).
A slight decrease in H v observed at the edges of the SPS-consolidated samples caused by the specific features of SPS processing. Due to the difference in electrical conductivity between the sintered material and electrically conducting graphite die, more electric current pulses pass through the die than through the sample. Thus, the sintering temperature at the edges is slightly higher, causing faster grain growth. Therefore, the H v at the edges is slightly smaller than in the center for all SPS-consolidated HEAs.
For the HPT-deformed bulk samples, on the contrary, a slight increase in hardness (Figure 11, blue) was observed from the center to edge due the expected larger imposed strain by torsion at the edge of the disk than in the center, which is typical for materials processed by high-pressure torsion [30]. SEM/EDX results of the HPT-deformed CoCrFeNi-