Effect of Ti on Characterization and Properties of CoCrFeNiTix High Entropy Alloy Prepared Via Electro-Deoxidization of the Metal Oxides and Vacuum Hot Pressing Sintering Process

The CoCrFeNi system is one of the most important high entropy alloys (HEAs) systems. By adding and adjusting the alloy element components and using different synthesis methods, different phases, organization and microstructure can be obtained, thus improving their properties. In this study, CoCrFeNiTix HEAs with various Ti contents (x in molar ratio, x = 0, 0.5, 1.0, 1.5) were fabricated by an electrochemical process by virtue of different oxides. The impacts of different Ti contents on the structure, distribution of elements, mechanical properties and corrosion behavior were researched using XRD, EDX and other testing methods. The bulk CoCrFeNiTix (x = 0, 0.5, 1.0, 1.5) HEAs could be obtained through vacuum hot pressing sintering process (VHPS), which had a single-phase FCC structure. The results of the study showed that the bulk CoCrFeNiTix exhibited superior ultimate tensile strength (UTS) and hardness, with the UTS of CoCrFeNiTi as high as 783 MPa and the hardness of CoCrFeNiTi1.5 reaching 669 HV. The corrosion behavior of CoCrFeNiTix (x = 0, 0.5, 1.0, 1.5) HEAs in 0.5 M H2SO4, 1 M KOH and 3.5 wt% NaCl was improved with addition of Ti. CoCrFeNiTix (x = 0, 0.5, 1.0, 1.5) HEAs have great potential for application in the fields of biomedical coating and aerospace, as well as extreme military industry, etc.


Introduction
High entropy alloys (HEAs) are a new group of engineering materials with multiple primary elements and interesting mechanical properties what have recently attracted attention in the field of metallic materials research. Their simple crystal structures offer a new way to develop advanced alloys that display tailored mechanical properties, unlike conventional alloys, which are based on one or two major elements [1][2][3][4][5][6][7]. Five or more elements compose high entropy alloys through atomic ratios that are the same or almost the same and which have concentrations of major elements ranging from 5at% to 35at%. The distinctive composition gives HEAs four core effects: sluggish diffusion effect, severe latticedistortion effect, high-entropy effect and cocktail effect [8]. HEAs tend to form stable solid solution phases and have superior properties, such as high strength, superior oxidation resistance, tensile strength, facture resistance and corrosion. FCC based CoCrFeNi alloys are classified as HEAs because of their high mixed entropy and excellent room-and lowtemperature mechanical properties, which have attracted significant academic attention and research [9][10][11][12][13].
Many researches have reported the impacts of various Ti concentration on phase, microstructure and properties of CoCrFeNiTi x HEAs. Research demonstrates that CoCrFeNiTi x HEAs changed from simple FCC structure to σ phase, R phase or Laves phase intermetallic compounds in HEAs with variation of Ti concentration [14][15][16]. CoCrFeNiTi x HEAs fabricated by different preparation methods, such as directed energy deposition high throughput synthesis [17], gas atomization [18], mechanical alloying [19] and electric arc melting [20] showed superior mechanical performance [21,22]. In the meantime, it has been demonstrated that post-treatments, such as aging [23,24], annealing [25] and thermomechanical processing [26,27] can enhance the mechanical performance of HEAs. In addition, there are numerous relevant reports on corrosion resistance [28], magnetic and electrical properties [19,29], and anti-irradiation resistance [30].
Many different methods for the preparation of HEAs have been proposed in recent years, such as: gas atomization [18], mechanical alloying [19] and electric arc melting [20]. The electric arc melting method is used most frequently, but it has the disadvantage of a high experimental temperature, which leads to significant energy consumption. Mechanical alloying is a solid-phase process, but materials produced by this method are susceptible to oxidation, and the cost of raw materials is too high. In the end, the bulk product needs to be prepared by sintering. The gas atomization method is complex and expensive, making it difficult to conduct in a laboratory.
Chen et al. propose the electrochemical approach [31] that has been studied in the fabrication of metal monomers and alloys [32][33][34]. In comparison with classical preparation processes, the fabrication of HEAs by electrochemical process has numerous advantages, such as low temperature, short process flow, and green environment. Electrochemical synthesis of HEAs, such as CoCrFeNi [35,36] and TiNbTaZr [37], via mixed oxides have been demonstrated, but the post-treatment and properties of these products have not been sufficiently studied.
In this paper, CoCrFeNiTi x HEAs powders of various Ti contents (x = 0, 0.5, 1.0, 1.5) were fabricated via electrochemical approach. Bulk CoCrFeNiTi x (x = 0, 0.5, 1.0, 1.5) HEAs were fabricated via VHPS process to investigate the mechanical and electrochemical corrosion properties. CoCrFeNiTi x HEAs are selected for the study because of their excellent mechanical properties.

Selection and Proportioning of Raw Materials
In this paper, we are using oxide electro-deoxygenation to prepare CoCrFeNiTi x (x = 0, 0.5, 1.0, 1.5) HEAs, so the mixed powders of CoO, Cr 2 O 3 , NiO, Fe 2 O 3 and TiO 2 were selected as raw materials based on their cost and ease of deoxygenation. Table 1 displays the composition of oxide powders for CoCrFeNiTi x HEAs obtained from the atomic ratio calculations.

Electro-Deoxidation Process of Products
The water-free CaCl 2 was laid in an alumina crucible and desiccated at 573 K for 12 h. The molten salt was pre-electrolyzed at 1173 K with nickel flake as the cathode and graphite as the anode at 2.8 V, and argon gas was continuously injected. Then, the oxide powders were electro-deoxygenated for 8 h at 3.1 V. The products were cleaned by the ultrasonic cleaner and dried under vacuum at 373 K.

Vacuum Hot Pressing Sintering
The products were ground and sieved. Then, the CoCrFeNiTi x (x = 0, 0.5, 1.0, 1.5) powders were loaded into graphite molds and hot pressed in a vacuum furnace. The temperature during VHPS was 1373 K, the rate of heating was 30 K/min, and the pressure was 35 MPa for 1 h. Bulk CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) HEAs samples were prepared and tested for their mechanical performance as well as corrosion resistance finally.

Characterization and Test
The electrolytic voltage was provided by a DC power supply (DP310, MESTEK, Shenzhen, China). The phase of bulk CoCrFeNiTi x high-entropy alloy samples with different Ti contents was examined by X-ray diffraction (D/max 2500PC, Rigaku, Japan). The microstructure, the fracture cross-section of bulk HEAs and elemental distribution of samples were characterized by scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDX) (TESCANVEGA II with Oxford INCA Energy 350). The remaining oxygen contents were measured via an oxygen nitrogen hydrogen analyzer (THC600, Germany). Carbon contents in alloys were analyzed using a carbon sulfur analyzer (G4 Icarus, Germany). The composition of powders was determined by inductively coupled plasma optical emission spectrometer (ICAP6300 DUO, ThermoFisher Scientific, Waltham, MA, USA).
Tensile specimens were polished and ground after hot pressing. Tensile performance was measured using a tensile testing machine with an original strain rate of 10 −4 s −1 . Two tests were performed for each sample, and good reproducibility was found. Hardness was measured with a load of 150 g and loading speed of 70 mm/s for 10 s via Vickers hardness tester (HV-115 type, Mitutoyo, Kanagawa, Japan) Corrosion resistance was measured via an electrochemical workstation. (CHI 660, Shanghai Chenhua Instrument Co. Ltd., Shanghai, China). The working electrodes were HEAs samples. The reference electrode was Ag/AgCl and the counter electrode was a platinum sheet. The scanning rate was 2 mV/s.

Electro-Deoxidization of the Mixed Oxides Powders
The chemical composition of the electro-deoxidization products is listed in Table 2 (at% is atomic fractions). Chemical composition is close to the set compositions of CoCrFeNiTi x high entropy alloy. There were two possible reasons for the difference between the electrodeoxidization products composition and the composition of the CoCrFeNiTi x high entropy alloy. First, some oxides had solubility in molten salts which caused the oxides to dissolve in the molten salt. Second, the fine particles produced in the process of electro-deoxidization might pass through the stainless steel mesh and lead to the loss of some oxides particles.  Figure 1 shows the SEM-BSE images and EDX analysis of the CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) HEAs powders. From low-magnification SEM-BSE images (see in Figure 1a-d), the products powders showed clusters composed of nodular particles. The higher magnification SEM-BSE images (see in Figure 1e-h) show that the nodular particle size of the products is bigger because of the increase in Ti contents. EDX composition maps (see in Figure 1i-l) confirmed that the distribution of elements was homogeneous at the micrometer level. 1, 1.5) HEAs powders. From low-magnification SEM-BSE images (see in Figure 1a-d), th products powders showed clusters composed of nodular particles. The higher magnifi cation SEM-BSE images (see in Figure 1e-h) show that the nodular particle size of the products is bigger because of the increase in Ti contents. EDX composition maps (see in Figure 1i-l)) confirmed that the distribution of elements was homogeneous at the mi crometer level.  Table 3 shows oxygen content and carbon content of electro-deoxidation product powders. The Ti content of CoCrFeNiTix (x = 0, 0.5, 1, 1.5) HEAs powders were 0 wt%, 0.5 wt%, 1 wt% and 1.5 wt%. It can be seen from Table 3 that the oxygen content of alloy powders is proportional to the Ti content. The formation of porous metals can be at tributed to how the removal of oxygen from solid oxides leaves oxygen vacancies, al lowing porous metals to be formed, which has been confirmed in the literature [38][39][40][41][42] This means that the molar volume of metal (Vm = Mm/ρm) should be smaller than th molar volume of its oxide (Vo = Mo/nρo) where m and o are the metal and oxide, re spectively, V is the molar volume, M is the molar mass, ρ is the density, and n is the number of metal atoms in the molecular formula of the oxides. The Vm/Vo ratio o Ti/TiO2 was close to 0.63, which was small. However, the Vm/Vo ratio of Ti/TiO was very close to 0.91. TiO was the intermediate phase in the late electro-reduction process of TiO2 Therefore, the dynamics of TiO reduction to Ti were difficult because the inherent po rosity of the TiO metal layer was small, especially considering the inevitabl high-temperature sintering of the metal. This caused an increase in the oxygen content o the electro-deoxidation products powders. As explained in the previous sections, ca thodic carbon deposition is caused by side reactions of CO3 2-and Ca 2+ . There were a number of factors affecting carbon deposition, mainly related to molten salt composition  Table 3 shows oxygen content and carbon content of electro-deoxidation products powders. The Ti content of CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) HEAs powders were 0 wt%, 0.5 wt%, 1 wt% and 1.5 wt%. It can be seen from Table 3 that the oxygen content of alloy powders is proportional to the Ti content. The formation of porous metals can be attributed to how the removal of oxygen from solid oxides leaves oxygen vacancies, allowing porous metals to be formed, which has been confirmed in the literature [38][39][40][41][42]. This means that the molar volume of metal (Vm = Mm/ρm) should be smaller than the molar volume of its oxide (Vo = Mo/nρo) where m and o are the metal and oxide, respectively, V is the molar volume, M is the molar mass, ρ is the density, and n is the number of metal atoms in the molecular formula of the oxides. The Vm/Vo ratio of Ti/TiO 2 was close to 0.63, which was small. However, the Vm/Vo ratio of Ti/TiO was very close to 0.91. TiO was the intermediate phase in the late electro-reduction process of TiO 2 . Therefore, the dynamics of TiO reduction to Ti were difficult because the inherent porosity of the TiO metal layer was small, especially considering the inevitable high-temperature sintering of the metal. This caused an increase in the oxygen content of the electro-deoxidation products powders. As explained in the previous sections, cathodic carbon deposition is caused by side reactions of CO 3 2and Ca 2+ . There were a number of factors affecting carbon deposition, mainly related to molten salt composition, electro-deoxidation time and reaction temperature. We therefore researched these factors in depth.

Structural and Morphological Characterization of the VHPS Product
The powders of electro-deoxidization products were hot pressed and sintered at 1373 K and 30 MPa for 1 h to obtain bulk CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) alloys. Figure 2 presents the XRD images of the bulk CoCrFeNiTi x alloys with different titanium contents. Three distinct XRD characteristic peaks were found near 2θ of 45 • , 50 • and 75 • , indicating that the main phase of CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) alloys was FCC phase solid solution. The higher mixing entropy of the alloys was the main reason for the formation of FCC solid solution. The presence of Laves phase (Co 2 Ti), R phase (Ni 3 Ti) and σ phase (FeCr) was observed in the ingot of CoCrFeNiTi x alloys prepared by the melting method, which has been confirmed in the literature [43][44][45]. The reason was related to the high temperature of the melting method and the elemental segregation during the solidification of the liquid metal in the melting process. The oxide precursor was deoxidized and alloyed during the electro-deoxidization process. There was no liquefaction of metals, and the process of alloying was dominated by solid phase diffusion, which prevented the formation of other intermetallic compound phases.

Structural and Morphological Characterization of the VHPS Product
The powders of electro-deoxidization products were hot pressed and sint 1373 K and 30 MPa for 1 h to obtain bulk CoCrFeNiTix (x = 0, 0.5, 1, 1.5) alloys. F presents the XRD images of the bulk CoCrFeNiTix alloys with different titanium co Three distinct XRD characteristic peaks were found near 2θ of 45°, 50° and 75°, ind that the main phase of CoCrFeNiTix (x = 0, 0.5, 1, 1.5) alloys was FCC phase solid s The higher mixing entropy of the alloys was the main reason for the formation solid solution. The presence of Laves phase (Co2Ti), R phase (Ni3Ti) and σ phase was observed in the ingot of CoCrFeNiTix alloys prepared by the melting method has been confirmed in the literature [43][44][45]. The reason was related to the high t ature of the melting method and the elemental segregation during the solidificatio liquid metal in the melting process. The oxide precursor was deoxidized and during the electro-deoxidization process. There was no liquefaction of metals, process of alloying was dominated by solid phase diffusion, which prevented mation of other intermetallic compound phases.  Figure 3 shows the SEM-BSE images and EDX analysis of the bulk CoCrFeN 0, 0.5, 1, 1.5) HEAs. From the low-magnification SEM-BSE images, darkly colo ferent phases dispersed in HEAs matrix. It can be seen from high-magnification SE images that, as the titanium increases, the size of the light phases gradually decrea the dark phases slowly increases. EDX analysis of the bulk CoCrFeNiTix HEAs  Figure 3 shows the SEM-BSE images and EDX analysis of the bulk CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) HEAs. From the low-magnification SEM-BSE images, darkly colored different phases dispersed in HEAs matrix. It can be seen from high-magnification SEM-BSE images that, as the titanium increases, the size of the light phases gradually decreases and the dark phases slowly increases. EDX analysis of the bulk CoCrFeNiTi x HEAs shows that the light phase is dominated by the segregation of Cr and the dark phase is dominated by titanium. The increase of titanium contents led to a gradual decrease in Cr segregation and a gradual increase in titanium segregation. In addition, the segregation size of titanium was small and presented particle dispersion, which was different from Cr segregation aggregation. Carbon deposition at the cathode was the main cause of Cr segregation. Comparison of CoCrFeNi and CoCrFeNiTi 0.5 revealed the same distribution of C and Cr elemental segregation, which could be explained by the formation of Cr 7 C 3 between deposited C and Cr during hot-press sintering. Comparison of CoCrFeNiTi and CoCrFeNiTi 1.5 revealed the same distribution of C and Ti segregation, which meant that TiC was formed between deposited C and Ti during hot-press sintering. With the increase of Ti, the C in the alloys changed from Cr 7 C 3 to TiC because C tended to form a TiC phase with Ti. In their experiments, Y.B. Peng et al. [46] found that the source of C could also be the graphite mold due to the high sintering temperature as well as the long VHC process; these factors can cause carbon contamination in the graphite mold.
that the light phase is dominated by the segregation of Cr and the dark phase is domi nated by titanium. The increase of titanium contents led to a gradual decrease in Cr seg regation and a gradual increase in titanium segregation. In addition, the segregation size of titanium was small and presented particle dispersion, which was different from C segregation aggregation. Carbon deposition at the cathode was the main cause of C segregation. Comparison of CoCrFeNi and CoCrFeNiTi0.5 revealed the same distribution of C and Cr elemental segregation, which could be explained by the formation of Cr7C between deposited C and Cr during hot-press sintering. Comparison of CoCrFeNiTi and CoCrFeNiTi1.5 revealed the same distribution of C and Ti segregation, which meant tha TiC was formed between deposited C and Ti during hot-press sintering. With the in crease of Ti, the C in the alloys changed from Cr7C3 to TiC because C tended to form a TiC phase with Ti. In their experiments, Y.B. Peng et al. [46] found that the source of C could also be the graphite mold due to the high sintering temperature as well as the long VHC process; these factors can cause carbon contamination in the graphite mold.  Figure 4 shows the stress-strain curves of the CoCrFeNiTix (x = 0, 0.5, 1, 1.5) HEAs. I can be observed from Figure 4 that the UTS of the CoCrFeNiTix HEAs is closely related to Ti. In general, increasing Ti contents was more beneficial for improving the UTS of the HEAs, but their relationship was not a monotonic increase. The UTS of the HEA changed from 670 MPa (Ti0) to 780 Mpa (Ti1). However, the UTS of the HEAs changed from 780 MPa (Ti1) to 460 MPa (Ti1.5). Therefore, the UTS of CoCrFeNiTix alloys showed a phenomenon of first increasing then later decreasing as the Ti content increased. Solid solution strengthening of the alloy was the main factor for the increase in UTS of the HEAs. The small amount of C deposited in the powder sample generates TiC with T during the hot pressing and sintering process. The phase stabilized TiC second-phase plasmas were diffusely distributed in the alloy matrix, which played the role of sec ond-phase strengthening and also contributed to the increase of UTS of CoCrFeNiTix al  Figure 4 shows the stress-strain curves of the CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) HEAs. It can be observed from Figure 4 that the UTS of the CoCrFeNiTi x HEAs is closely related to Ti. In general, increasing Ti contents was more beneficial for improving the UTS of the HEAs, but their relationship was not a monotonic increase. The UTS of the HEAs changed from 670 MPa (Ti 0 ) to 780 Mpa (Ti 1 ). However, the UTS of the HEAs changed from 780 MPa (Ti 1 ) to 460 MPa (Ti 1.5 ). Therefore, the UTS of CoCrFeNiTi x alloys showed a phenomenon of first increasing then later decreasing as the Ti content increased. Solid solution strengthening of the alloy was the main factor for the increase in UTS of the HEAs. The small amount of C deposited in the powder sample generates TiC with Ti during the hot pressing and sintering process. The phase stabilized TiC second-phase plasmas were diffusely distributed in the alloy matrix, which played the role of second-phase strengthening and also contributed to the increase of UTS of CoCrFeNiTi x alloy. The elongation of the CoCrFeNiTi x alloy also showed a phenomenon of first increasing (Ti 0~T i 0.5 alloy) then later decreasing (Ti 0.5~T i 1.5 alloy) as the Ti content increased. Ti 0.5 alloy had excellent elongation at break compared to other CoCrFeNiTi x (x = 0, 0.5, 1, 1.5) HEAs.

Tensile and Hardness Results
loy. The elongation of the CoCrFeNiTix alloy also showed a phenomenon of first increasing (Ti0~Ti0.5 alloy) then later decreasing (Ti0.5~Ti1.5 alloy) as the Ti content increased. Ti0.5 alloy had excellent elongation at break compared to other CoCrFeNiTix (x = 0, 0.5, 1, 1.5) HEAs. As is commonly known, the Ti-O chemical bond was stable; this made the deep deoxygenation of TiO2 exceptionally difficult and resulted in an increase of the residual unremoved oxygen content in the product alloy powder as the Ti content increase. The residual oxygen formed a secondary TCP phase in the matrix. Stress concentration and cracking were induced during the tensile process, and the reason for the significant decrease in alloy toughness was the accumulation of interfacial dislocations caused by the large asymmetry between the TCP phase and the FCC phase [47]. Figure 5 presents the fracture surfaces of CoCrFeNiTix HEAs. The dimple morphology was exhibited in Ti0 and Ti0.5 HEAs (see Figure 5a,b) because the FCC phase was the main physical phase of the CoCrFeNiTix HEAs; this caused it to exhibit significant toughness fracture. The increased oxygen contents of the HEAs caused the Ti1.0 alloy and Ti1.5 alloy samples to show river pattern detachment fracture morphology that is typical of brittle fracture.  Figure 6 shows the hardness of the tensile test samples. The hardness of CoCrFeN-iTix HEAs gradually increased as Ti increased, from 245 HV in Ti0 alloy to 669 HV in Ti1.5 alloy. There were two explanations for this phenomenon: first, the Ti atoms with larger atomic radius caused lattice distortion in the high-entropy alloy, and the lattice distortion became more serious because of the increase of Ti contents [48]; second, carbon and Ti As is commonly known, the Ti-O chemical bond was stable; this made the deep deoxygenation of TiO 2 exceptionally difficult and resulted in an increase of the residual unremoved oxygen content in the product alloy powder as the Ti content increase. The residual oxygen formed a secondary TCP phase in the matrix. Stress concentration and cracking were induced during the tensile process, and the reason for the significant decrease in alloy toughness was the accumulation of interfacial dislocations caused by the large asymmetry between the TCP phase and the FCC phase [47]. Figure 5 presents the fracture surfaces of CoCrFeNiTi x HEAs. The dimple morphology was exhibited in Ti 0 and Ti 0.5 HEAs (see Figure 5a,b) because the FCC phase was the main physical phase of the CoCrFeNiTi x HEAs; this caused it to exhibit significant toughness fracture. The increased oxygen contents of the HEAs caused the Ti 1.0 alloy and Ti 1.5 alloy samples to show river pattern detachment fracture morphology that is typical of brittle fracture.
loy. The elongation of the CoCrFeNiTix alloy also showed a phenomenon of first increasing (Ti0~Ti0.5 alloy) then later decreasing (Ti0.5~Ti1.5 alloy) as the Ti content increased. Ti0.5 alloy had excellent elongation at break compared to other CoCrFeNiTix (x = 0, 0.5, 1, 1.5) HEAs. As is commonly known, the Ti-O chemical bond was stable; this made the deep deoxygenation of TiO2 exceptionally difficult and resulted in an increase of the residual unremoved oxygen content in the product alloy powder as the Ti content increase. The residual oxygen formed a secondary TCP phase in the matrix. Stress concentration and cracking were induced during the tensile process, and the reason for the significant decrease in alloy toughness was the accumulation of interfacial dislocations caused by the large asymmetry between the TCP phase and the FCC phase [47]. Figure 5 presents the fracture surfaces of CoCrFeNiTix HEAs. The dimple morphology was exhibited in Ti0 and Ti0.5 HEAs (see Figure 5a,b) because the FCC phase was the main physical phase of the CoCrFeNiTix HEAs; this caused it to exhibit significant toughness fracture. The increased oxygen contents of the HEAs caused the Ti1.0 alloy and Ti1.5 alloy samples to show river pattern detachment fracture morphology that is typical of brittle fracture.  Figure 6 shows the hardness of the tensile test samples. The hardness of CoCrFeN-iTix HEAs gradually increased as Ti increased, from 245 HV in Ti0 alloy to 669 HV in Ti1.5 alloy. There were two explanations for this phenomenon: first, the Ti atoms with larger atomic radius caused lattice distortion in the high-entropy alloy, and the lattice distortion became more serious because of the increase of Ti contents [48]; second, carbon and Ti  Figure 6 shows the hardness of the tensile test samples. The hardness of CoCrFeNiTi x HEAs gradually increased as Ti increased, from 245 HV in Ti 0 alloy to 669 HV in Ti 1.5 alloy. There were two explanations for this phenomenon: first, the Ti atoms with larger atomic radius caused lattice distortion in the high-entropy alloy, and the lattice distortion became more serious because of the increase of Ti contents [48]; second, carbon and Ti generate hard TiC second-phase particles, and the hardness of the alloys was significantly increased by the strengthening effect of second-phase particles. generate hard TiC second-phase particles, and the hardness of the alloys was significantly increased by the strengthening effect of second-phase particles.  Figure 7a that a passive region was formed in the total curves range of -0.3-0.9 VSHE. All HEAs samples showed a secondary passive region, and the difference in the range of passive region of the HEAs samples was not significant. It can be seen from Table 4 that Ti can significantly improve the Ecorr of the HEAs, but the increase of Epit is not obvious. The corrosion current density was lower than 304 stainless steel and pure titanium at 0.5 M H2SO4 [49], indicating that the addition of Ti could promote the formation of corrosion-resistant high-performance oxide films. Ti could easily form an oxide film on metal surfaces in 0.5 M H2SO4. The oxide film had excellent acid corrosion resistance, and the thickness had little significant effect on corrosion resistance. Figure 7b shows the stable polarization curve of CoCrFeNiTix alloy at 1M KOH. Figure 7b shows that, on the surface of the HEA, a clear passive region was formed between the curves and a stable oxide film was formed in 1 M KOH. The passivated region range of Ti0 alloy was significantly wider than the passive region of Ti0.5, Ti1.0 and Ti1.5 alloys. As can be seen from Table 5, the Ti0 alloy had high Epit, and the addition of Ti did not obviously improve the Ecorr of the alloy, and the icorr was slightly reduced; these items indicate that the corrosion resistance of the HEAs improved slightly with the addition of Ti in 1 M KOH. Figure 7c shows the stable polarization curve of CoCrFeNiTix alloy at 3.5 wt% NaCl, and that the passive region range of Ti1.5 alloy is significantly narrower than the passive region of Ti0, Ti0.5 and Ti1.0 alloys. This might be because the Ti1.5 alloy with higher oxygen content was more prone to localized and pitting corrosion in the oxygen enriched region during polarization tests, where the passivated film was struck through and the extent of the passive region was reduced. The increase in Ti contents leads to an increase in the nano-dispersion of TiN, resulting in the formation of a passive film of TiO2 during the corrosion process. The breakage potential of passivation in pitting corrosion increases with the increase of TiN contents in the film. [50] Table 6 shows that the Ecorr of the HEAs ranged from −1.138 to −1.224 VSCE and the Epit ranged from −0.142 to −0.512 VSCE. Ti0 and Ti0.5 HEAs had higher Ecorr and Epit because the oxide precursor deoxidization was more difficult with the increase of Ti, and the elevated oxygen content in the product reduced It can be observed from Figure 7a that a passive region was formed in the total curves range of −0.3-0.9 V SHE . All HEAs samples showed a secondary passive region, and the difference in the range of passive region of the HEAs samples was not significant. It can be seen from Table 4 that Ti can significantly improve the E corr of the HEAs, but the increase of E pit is not obvious. The corrosion current density was lower than 304 stainless steel and pure titanium at 0.5 M H 2 SO 4 [49], indicating that the addition of Ti could promote the formation of corrosion-resistant high-performance oxide films. Ti could easily form an oxide film on metal surfaces in 0.5 M H 2 SO 4 . The oxide film had excellent acid corrosion resistance, and the thickness had little significant effect on corrosion resistance.      Figure 7b shows the stable polarization curve of CoCrFeNiTi x alloy at 1M KOH. Figure 7b shows that, on the surface of the HEA, a clear passive region was formed between the curves and a stable oxide film was formed in 1 M KOH. The passivated region range of Ti 0 alloy was significantly wider than the passive region of Ti 0.5 , Ti 1.0 and Ti 1.5 alloys. As can be seen from Table 5, the Ti 0 alloy had high E pit , and the addition of Ti did not obviously improve the E corr of the alloy, and the i corr was slightly reduced; these items indicate that the corrosion resistance of the HEAs improved slightly with the addition of Ti in 1 M KOH.  Figure 7c shows the stable polarization curve of CoCrFeNiTi x alloy at 3.5 wt% NaCl, and that the passive region range of Ti 1.5 alloy is significantly narrower than the passive region of Ti 0 , Ti 0.5 and Ti 1.0 alloys. This might be because the Ti 1.5 alloy with higher oxygen content was more prone to localized and pitting corrosion in the oxygen enriched region during polarization tests, where the passivated film was struck through and the extent of the passive region was reduced. The increase in Ti contents leads to an increase in the nano-dispersion of TiN, resulting in the formation of a passive film of TiO 2 during the corrosion process. The breakage potential of passivation in pitting corrosion increases with the increase of TiN contents in the film [50]. Table 6 shows that the E corr of the HEAs ranged from −1.138 to −1.224 V SCE and the E pit ranged from −0.142 to −0.512 V SCE . Ti 0 and Ti 0.5 HEAs had higher E corr and E pit because the oxide precursor deoxidization was more difficult with the increase of Ti, and the elevated oxygen content in the product reduced the corrosion resistance of the HEAs. Secondly, the secondary phase in the high Ti alloy also deteriorated the corrosion resistance of the HEAs.

Conclusions
The CoCrFeNiTi x HEAs with different Ti contents were prepared by electro-deoxidization of metal oxides. The products powders showed clusters composed of nodular particles, and the distribution of elements was homogeneous. The oxygen content of alloys gradually increased as Ti increased. The CoCrFeNiTi x HEAs had an FCC structure after hot pressing. The increase of Ti significantly improved Cr elemental segregation in alloys. With the increase of Ti, UTS of the CoCrFeNiTi x HEAs first increased (Ti 0 -Ti 1.0 ) and then decreased (Ti 1.0 -Ti 1.5 ). The UTS of CoCrFeNiTi was up to 783 MPa, and the hardness of alloys gradually increased as Ti increased, reaching up to 669 HV. The corrosion resistance results displayed that the addition of Ti significantly improved the acid corrosion resistance of the alloy in 0.5 M H 2 SO 4 solution; the corrosion resistance of alloys slightly increased with addition of Ti in 1 M KOH solution; Ti 0 and Ti 0.5 alloys had higher corrosion potential and pitting potential in 3.5 wt% NaCl solution; and Ti 0.5 alloy had the best corrosion resistance.