Influence of Oxide Dispersions (Al₂O₃, TiO₂, and Y₂O₃) in CrFeCuMnNi High-Entropy Alloy on Microstructural Changes and Corrosion Resistance

Abstract

This study investigates the influence of 3 vol.% Al₂O₃, 3 vol.% TiO₂, and 3 vol.% Y₂O₃ in the CrFeCuMnNi equimolar high-entropy alloy on its microstructural changes and corrosion resistance. These oxide-dispersed high-entropy composites (ODS-HECs) were synthesized via high-energy ball milling (50 h) followed by uniaxial hot-compaction (550 MPa, 45 min), medium-frequency sintering (1100 °C, 20 min), and hot forging (50 MPa). The microstructures of the developed composites produced a stable FCC phase, a small amount of ordered BCC-B2 structure, Fe₂O₃, and corresponding dispersed oxide phases. The corrosion of the developed high-entropy composites was tested in 3.5% NaCl solution using several electrochemical techniques. The results revealed that the corrosion rate (R_Corr) decreased with the incorporation of oxide particles. Among the investigated samples and based on the electrochemical impedance spectroscopy results, CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs were seen to possess the highest value of corrosion resistance (R_P). The change in the chronoamperometric current with time indicated that the CrFeCuMnNi alloy suffered pitting corrosion which decreased when Al₂O₃ was added, forming a CrFeCuMnNi-3 vol.% Al₂O₃ sample. In contrast, the incorporation of a 3 vol.% Y₂O₃, and 3 vol. TiO₂, prevents pitting.

1. Introduction

Globally, the development of new materials for high-temperature applications, such as materials for aerospace parts, nuclear reactors, military components, oil industrial parts, automotive components, and structural components, is necessary and important in the current scenario [1]. Ferritic and martensitic steels have been considered as potential materials for defense applications in the past decades because of their excellent thermal stability, high thermal conductivity, low neutron irradiation effect, and excellent corrosion resistance [2]. Furthermore, to enhance the product swelling resistance at high temperatures (to achieve almost zero), and enrich the temperature-withstanding properties (up to 800 °C at least), oxide ceramic particles were incorporated into ferritic and martensitic steels to obtain oxide-dispersion-strengthened (ODS) steels [3]. However, the expected level was not achieved owing to the increased weight and moderate corrosion at elevated temperatures.

In general, the properties of metals and alloys can be completely changed by structural refinements in terms of dislocations, lattice distortion, the formation of a greater number of nano-grains, an increase in the surface energy, and an effective interface between the elements and atoms. Structural refinement can be easily achieved via high-energy mechanical alloying (MA), which is one of the solid-state powder processing techniques. This MA process involves the severe plastic deformation of charged metallic materials, the cold welding and fracturing of powder particles, and repeated mechanical collisions to introduce structural refinements that alter the physical, electrical, optical, chemical, and mechanical properties [4]. Concentrated solid solution alloys (CSAs) can provide extensive microstructural features and remarkable properties, among which high-entropy alloys (HEAs) are a type of CSA. Normally, CSAs have a single-phase body-centered cubic (BCC) structure or a face-centered cubic (FCC) structure, a hexagonally close-packed (HCP) structure, or a mixture of these phases [5,6,7].

Recently, HEAs have become one of the new categories of engineering alloys which are currently being investigated by several researchers [8]. HEAs are multicomponent systems in which more than five metallic elements are atomically mixed to achieve various properties. However, the phase stability at elevated temperatures is challenging because of grain coarsening at high temperatures [9]. Nano-structural formations in HEAs along with high configurational entropy have shown improved mechanical properties and wear resistance [10,11,12]. Vaidya et al. [11] synthesized CoCrFeNi and CoCrFeNiMn HEAs via the MA route. The HEA powders were consolidated via spark plasma sintering (SPS) and the authors found that enhanced strength with stable phases was achieved. Wang et al. [13] manufactured Ni_1.5CoCuFeCr_0.25V_0.25 and Ni_1.5CoCuFeV_0.5 HEAs using MA followed by SPS, and improved results were obtained. Gludovatz et al. [14] synthesized and developed a CoCrFeNiMn HEA for cryogenic applications. This alloy exhibited high fracture toughness values of 0.219 GPam^−1/2 at −196 °C, 0.221 GPam^−1/2 at −93 °C, and 0.217 GPam^−1/2 at 0 °C. However, this HEA exhibited distorted tensile curves owing to its structural complexity and pinning of dislocations at a lower temperature. Hence, to retain the mechanical, wear, corrosion, and thermal properties at high temperatures, the concept of oxide dispersion strengthening (ODS) over single-phase or binary phases, or multiphases of CSAs or HEAs, is currently of interest to materials scientists. This concept can refine the grains and pin down the grain growth in alloys (binary alloys, CSAs, and HEAs) [15].

The Introduction of oxide particles (Al₂O₃-alumina, TiO₂-titania, ThO₂-thoria, and Y₂O₃-yttria) into CSAs or HEAs is expected to improve their properties. The incorporation of these oxide particles into HEAs results in a greater number of interfaces, dislocation loops, and bubbles/voids among the grain boundaries (GBs), thereby enhancing the hardening effect [16]. In addition, the incorporation of oxide particles into CSAs or HEAs can provide advantages for both ODS alloys and HEAs. Furthermore, the phase stability and oxidation resistance of CSAs and HEAs can be improved by adding oxide particles [17]. Hadraba et al. [18] synthesized CoCrFeNiMn reinforced with 0.3 wt.% Y₂O₃ ODS-HEAs using the in situ reaction method. This material was developed by mixing all elemental powders in a Fritch ball mill, with a ball-to-powder ratio of 15:1 at 350 rpm for 24 h under vacuum. Oxygen (O₂), yttrium (Y), and titanium (Ti) were incorporated and MA was carried out to form Y₂O₃ in the alloy. The synthesized powders were consolidated via SPS at 1423 K under a pressure of 50 MPa. The microstructure and mechanical behavior were investigated and the mechanical results at room temperature revealed that the ultimate compressive strengths of the HEA and ODS-HEA were 1.010 GPa and 1.138 GPa, respectively. In addition, 50% refinement of grains was achieved in ODs-HEA compared to HEA.

Gwalani et al. [19] synthesized Al_0.3CoCrFeMnNi dispersed in 3 vol.% of yttria ODS-HEA through MA followed by SPS. The results revealed that the developed ODS-HEA exhibited 1.8 GPa strength which was 1.8 times higher compared to that of the HEA, owing to the dispersion strengthening produced by yttria particles. Liu et al. [20] developed a Cr-Mn-Fe-Co-Ni HEA dispersed with yttria (0.25 wt.%) prepared via MA and SPS. The metal powders related to the composition were added inside the ball mill, MA was carried out for 50 h at 300 rpm with a BPR of 10:1, and a stainless steel vial with zirconia balls was used as the grinding medium. The synthesized ODS-HEA powders were consolidated via SPS at 900 °C for 5 min at a pressure of 50 MPa. Phase evaluations and microstructural examinations were performed using XRD and advanced electron microscopy, respectively. In addition, the mechanical behavior and wear behavior were investigated. A tensile strength of approximately 1 GPa was achieved in a 0.25 wt.% yttria-incorporated HEA and it produced excellent wear resistance. Jia et al. [21] synthesized an FeCrCoNi-based HEA dispersed in 5 wt.% yttria via MA and SPS. The elemental powders related to the composition (more than 99.5% purity) were mixed in a high-energy ball mill, milled up to 60 h, 30 nm yttria was incorporated into the alloy, and the synthesized powders were consolidated via SPS at 1273 K for 10 min. The microstructural evolution and mechanical behavior were investigated in detail. The microstructural investigation indicated that the developed ODS-HEA sample produced a stable FCC phase and exhibited a tensile strength of approximately 1.75 GPa due to the dispersion of nano-sized yttria particles.

Yang et al. [22] synthesized Al_0.4FeCrCo_1.5NiTi_0.3 dispersed with nano alumina of ODS-HEA via MA. The synthesized samples produced a supersaturated solid solution of a single-phase FCC structure, consolidated to the bulk sample via SPS at 1273 k for 10 min at a pressure of 30 MPa, with a tensile strength of 2.14 GPa and Vickers hardness strength of approximately 14 GPa. Based on the literature, the addition of oxide particles effectively improves the phase stability, oxidation resistance, and mechanical properties of CSA/HEAs. There are no studies related to the development and investigation of CrFeCuMnNi HEA reinforced/dispersed with Y₂O₃, TiO₂, and Al₂O₃ of oxide-dispersed high-entropy composites (ODS-HECs); hence, the present research work was carried out. The main objectives of this study are as follows: synthesize CrFeCuMnNi, CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% TiO₂, and CrFeCuMnNi-3 vol.% Y₂O₃ via MA followed by hot compaction, medium-frequency sintering, and hot forging; examine the phase and microstructural formations; and investigate the corrosion performance using various electrochemical techniques such as potentiodynamic polarization, electrochemical impedance spectroscopy, and the variation in the chronoamperometric current for a certain time.

2. Materials and Methods

2.1. Synthesis and Consolidation of ODS-HECs

Pure elemental powders (more than 99% purity), namely, chromium (Cr, BCC), iron (Fe, BCC), copper (Cu, FCC), manganese (Mn, FCC), and nickel (Ni, FCC), were purchased from M/s Nanografi, Germany. Similarly, the oxide-dispersed particles of alumina (α-Al₂O₃, rhombohedral), titania (TiO₂, tetragonal), and yttrium oxide (Y₂O₃, cubic) were purchased from M/s Nanografi, Germany, with more than 99.5% purity.

Table 1. Chemical composition of synthesized oxide-dispersive strengthened high-entropy composites (ODS-HECs) of CrFeCuMnNi, CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% TiO₂, and CrFeCuMnNi-3 vol.% Y₂O₃.

Name of Sample (ID)	Cr		Fe		Cu		Mn		Ni		Reinforcements
	at.%	wt.%	at.%	wt.%	at.%	wt.%	at.%	wt.%	at.%	wt.%	at.%	wt.%
CrFeCuMnNi	0.200	0.196	0.200	0.182	0.200	0.223	0.200	0.193	0.200	0.206	0.000	0.000
CrFeCuMnNi-3 A2O3	0.194	0.186	0.194	0.173	0.194	0.211	0.194	0.186	0.194	0.195	0.030	0.052
CrFeCuMnNi-3 TiO2	0.194	0.188	0.194	0.175	0.194	0.214	0.194	0.185	0.194	0.197	0.030	0.042
CrFeCuMnNi-3 Y2O3	0.194	0.175	0.194	0.162	0.194	0.198	0.194	0.172	0.194	0.183	0.030	0.109

lists the chemical compositions of the synthesized ODS-HECs according to atomic and weight percentages. The elemental powder was weighed using an electronic balance and charged into a four-station high-energy ball mill (M/s. TENCAN, XQM-4A, Changsha City, Hunan Province, China). A ball-to-powder ratio of 15:1, milling speed of 350 rpm, and total milling time of 50 h were used for the synthesis of the powders. Milling was carried out using ethanol as a process control agent (PCA, wet medium) to minimize the cold welding and oxidization of powder particles. Sixty grams of powder in each vial was taken during ball milling, and intermittent milling of 15 min forward, 15 min pass, and 15 min reverse was used to avoid the over-heating of charged particles. Hardened stainless steel vials and balls were used as the grinding medium. The powders milled for 50 h were dried and stress was relieved under vacuum at a temperature of 120 °C for 30 min inside a tube furnace (M/s. Nabatherm, Lilienthal, Germany). Stress-relieved powders were then charged inside an H13 steel die-set, heated to 550 °C at a heating rate of 10 °C/min, held for 45 min, and then hot-compacted at 550 MPa and held for 10 min. Hot-compacted pellets with a diameter of 15 mm and height of 25 mm were sintered at 1200 ± 20 °C under argon gas using a medium-frequency electric induction furnace (M/s Zhengzhou Yuanjie Chemical Co., Ltd., Zhengzhou, China). A heat input of 2000 watts with a heating rate of 120°/min was set. Sintered hot-pellets were immediately forged at 1000 ± 40 °C with 50 MPa pressure, and cooled to room temperature in air. Figure 1 shows a schematic diagram representing the atomic structure of the as-received powders, synthesized ODS-HEC powders, hot compaction, medium-frequency sintering, and hot forging of the present research study.

2.2. Electrochemical Experiments

Hot-forged samples with a diameter of 15 mm and height of 10 mm were prepared for electrochemical measurements. All samples were ground using SiC papers with different grit sizes, and polished using alumina to ensure that the scratch was free with a mirror appearance. Sodium chloride solution (NaCl, 3.5%) received from Sigma Aldrich (99.9% purity) was used in electrochemical experiments which were selected based on the marine environment. The data obtained from the electrochemical instrument were collected using an Autolab Potentiostat Galvanostat (Ecochemie PGSTAT 30, Metrohm, Amsterdam, the Netherlands). These measurements were performed using a three-electrode cell that could accommodate 300 mL of the test solution. The samples investigated were used as the working electrodes (WEs). The samples were welded and mounted in inert epoxy. The full preparation of the working electrodes for electrochemical measurements was prepared as reported in our previous work [23]. A silver/silver chloride (Ag/AgCl) electrode was used as the reference electrode (RE). A platinum sheet was used as the counter electrode (CE). The cyclic potentiodynamic polarization (CPP) data were obtained by scanning the potential in the forward direction between −0.8 V and 0.2 V at 0.00166 V/s. The potential was rescanned in the backward direction at the same scan rate until the reversed current values intersected with the forward current values. Electrochemical impedance spectroscopy (EIS) plots were collected by applying a sinusoidal wave perturbation amplitude of ±5 mV for a frequency scan range between 100,000 Hz and 0.1 Hz from the values of the corrosion potential (OCP). A chronoamperometric current time (CCT) experiment was performed after applying two different potential values of −0.250 V and −0.050 V (Ag/AgCl) on the sample surface for 40 min. All of the electrochemical experiments were performed after immersing the samples in a chloride solution for 60 min before the measurements. Figure 2 shows a schematic of the electrochemical test used in this study.

2.3. Materials Characterization

The received metallic powders’ surface morphologies were examined using an Apreo field emission gun (FEG) high-resolution scanning electron microscope (HRSEM) operated at 30 keV, and the machine had 1.3 nm resolution at 1 keV. The same HRSEM technique was used to examine the microstructure of the hot-forged samples. An X-ray diffractometer made of M/s Empyrean, Malvern Panalytical, was used to examine the X-ray diffraction (XRD) profile of the as-received elemental powders (to check the purity), 50 h ball milled powders, and hot-forged samples. The XRD test was conducted at a scanning speed of 0.6°/min with a step size of 0.01°. X’Pert high score plus software was used to investigate the formation of various phases. After the corrosion test, the surface of the sample was examined using an HRSEM.

3. Results and Discussion

3.1. HRSEM and XRD Analyses on As-Received Elemental Powders

The surface morphology of the as-received elemental powder was examined using an HRSEM which is shown in Figure 3. Irregular flake/platelet powder particles with an average particle size of 48 ± 3.4 μm were observed in Cr (BCC, Figure 3a). Irregular and almost equiaxed powder particles with an average particle size of 45 ± 2.8 μm were observed in Fe (BCC, Figure 3b). Regular spherical powder particles with an average size of approximately 26 ± 6.45 μm were observed in the as-received Cu (FCC, Figure 3c). An irregular polygonal shape with an average size of 27 ± 5.78 μm was observed for Mn (FCC, Figure 3d). Almost spherical and satellite particles with an average size of around 20.5 ± 3.89 μm was observed in Ni (FCC, Figure 3e). No foreign particles were observed, which ensured the purity under the as-received condition. Figure 4 shows the HRSEM images (left) and the corresponding XRD peak profiles (right) of the as-received reinforcement particles of α-Al₂O₃, TiO₂, and Y₂O₃. The average agglomerate particle sizes of the as-received reinforcements were 105 nm, 90 nm, and 8 μm for α-Al₂O₃, TiO₂, and Y₂O₃, respectively. The crystal structures of the as-received reinforcements were examined using XRD peak profiles which were rhombohedral, tetragonal, and cubic for α-Al₂O₃, TiO₂, and Y₂O₃, respectively.

Figure 5a,b shows the HRSEM powder surface morphology of 50 h milled CrFeCuMnNi high-entropy alloy powders as an example, which is presented for ensuring the alloy formation. Very fine powder particles with some flake-like-shape morphology were obtained after high-energy ball milling (50 h). The average particle size of around 4.8 ± 1.6 μm was observed and all powder particles exhibited equiaxed morphology, indicating the steady-state attainment after 50 h MA [24]. In general, cold welding and fracturing are the repeated mechanisms involved in the MA process, leading to the production of very fine powder particles produced by severe plastic deformation. Figure 5b shows the magnified view of Figure 5a, indicating a bright image which confirms the alloy formation and homogeneous state. Here, the charged powder particles were subjected to being broken due to repeated plastic deformation obtained from the combination of shear and impact forces leading to the production of a single-phase solid solution as seen in Figure 5b [25]. An HRSEM-EDAX spectrum was also carried out (Figure 5c) which exhibits all peaks related to the incorporated metallic elements. The corresponding elemental composition is also given in Figure 5d.

3.2. XRD and HRSEM Phase Analyses on ODS-HECs

The X-ray peak profiles of the blended (0 h) CrFeCuMnNi, CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% TiO₂, and CrFeCuMnNi-3 vol.% Y₂O₃ are shown in Figure 6. All of the samples exhibited very sharp and narrow peaks with high-intensity values owing to the absence of energy imported in to the charged materials. Figure 7a shows the XRD peak profile of the 50 h mechanically alloyed oxide-dispersed high-entropy composites (ODS-HECs). Both Figure 6 and Figure 7a were drawn on the same scale. From Figure 7a, it can be seen that the peak intensity of the incorporated elements is drastically reduced and more peak broadening is observed after 50 h of MA. Several peaks disappeared and merged to form a single peak, confirming the formation of a solid solution. Powder samples milled for 50 h produced major BCC and FCC phases in addition to the incorporated reinforcement particle peaks (Al₂O₃, TiO₂, and Y₂O₃). These results confirmed the successful synthesis of ODS-HECs. The atomic radii of the incorporated metallic elements (Cr, Fe, Cu, Mn, and Ni) were almost identical (

Table 2. Elemental parameters of developed ODS-HEACs of CrFeCuMnNi matrix alloy.

Name of Element	Atomic Radius, nm	Crystal Structure	Melting Point (°C)	Valence Electron Concentration (VEC)	Theoretical Density, g/cm3
Cr	0.140	BCC	1538	8	7.140
Fe	0.140	BCC	1907	6	7.874
Cu	0.135	FCC	1084	11	8.920
Mn	0.140	FCC	1246	7	7.260
Ni	0.135	FCC	1455	10	8.908
α-Al2O3	-	Rhombohedral	2072	-	3.987
TiO2	-	Tetragonal	1843	-	4.260
Y2O3	-	Cubic	2425	-	5.010

), and hence a super saturated solid solution was easily obtained (Figure 5b). Drastic decreases in the peak intensity and peak broadening were attributed to the longer milling time, repeated cold welding, and the fracturing of powder particles produced by the kinetic energy of the input medium. Among the synthesized ODS-HEC powders, the CrFeCuMnNi-3 vol.% TiO₂ sample exhibited a very high value of peak broadening, indicating that more structural refinements occurred. Following this, the CrFeCuMnNi-3 vol.% Al₂O₃ sample produced a little decrease in peak broadening, and the CrFeCuMnNi-3 vol.% Y₂O₃ sample exhibited low peak broadening compared to all of the samples (Figure 7b). Figure 7b shows that peak shifting occurred in the CrFeCuMnNi HEA matrix owing to the addition of ceramic particles, confirming the structural refinement. The melting point of TiO₂ is low (1843 °C, Table 2), and hence TiO₂ can easily mix and adhere to the CrFeCuMnNi matrix, leading to more structural refinement [26]. However, the melting point of Y₂O₃ is high (2425 °C, Table 2) which is expected to offer more resistance for adhering to the CrFeCuMnNi matrix, causing less structural refinement [27]. These results indicate that little structural refinement occurred in Y₂O₃-reinforced samples.

Figure 8 shows the XRD peak profile of the hot-forged samples which exhibit a major/single FCC phase. The XRD results of the powder samples milled for 50 h demonstrate a major BCC phase (Figure 7) in all of the samples, which diminished or almost disappeared after hot forging. Here, high-temperature sintering at a medium frequency and hot forging followed by air cooling promoted the formation of the FCC phase. In addition to the major FCC phase, an ordered BCC-B2 phase was formed in all of the samples. The formation of FCC and ordered BCC-B2 phases has also been observed by other researchers [28,29]. The formation of FCC and ordered BCC-B2 phases is expected to improve the corrosion and mechanical performance of the developed ODS-HECs [30,31]. Furthermore, all of the hot-forged samples produced the Fe₂O₃ phase, which was expected to be formed from contaminants obtained from the grinding medium and high-temperature sintering.

Figure 9 shows the HR-SEM microstructures of the hot-forged samples of the developed ODS-HECs. The SEM microstructures of all of the forged samples showed major FCC, as well as a small amount of ordered BCC-B2, Fe₂O₃, and the corresponding reinforcement phases, as observed via the XRD results (Figure 8). A stable FCC structure was also observed by Jia et al. [21] while developing an FeCrCoNi HEA incorporated with Y₂O₃. The average grain size of the CrFeCuMnNi matrix was calculated using ImageJ software, which was 15.5 μm, 9.4 μm, 5.5 μm, and 12.6 μm for CrFeCuMnNi, CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% TiO₂, and CrFeCuMnNi-3 vol.% Y₂O₃, respectively. A small matrix grain size was observed in the CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs sample which was attributed to more structural refinement, less grain growth, and the uniform dispersion of nano-TiO₂ particles. However, the large matrix grain sizes in the unreinforced and CrFeCuMnNi-3 vol.% Y₂O₃ samples were expected to have less structural refinement, and more grain growth occurred in these samples. Structural refinement with the incorporation of Y₂O₃ in the CoCrFeNiMn HEA was also observed by Hadraba et al. [18]. These SEM results confirmed the successful fabrication of the ODS-HECs in this study.

3.3. Electrochemical Corrosion Behavior of ODS-HECs

3.3.1. Cyclic Potentiodynamic Polarization (CPP) Data

The CPP measurements were carried out to report the effects of Al₂O₃, Y₂O₃, and TiO₂ on the corrosion passivation of the CrFeCuMnNi high-entropy alloy matrix in 3.5% NaCl solution. Figure 10 displays the CPP curves obtained for the (a) CrFeCuMnNi, (b) CrFeCuMnNi-3 vol.% Al₂O₃, (c) CrFeCuMnNi-3 vol.% Y₂O₃, and (d) CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs after their immersion for 1 h in 3.5% NaCl solution. Corrosion parameters such as cathodic Tafel (βc) and anodic Tafel (βa) slopes, corrosion potential (E_Corr), corrosion current (j_Corr), protection potential (E_Prot), pitting potential (E_Pit), polarization resistance (R_P), and corrosion rate (R_Corr) were obtained and are listed in

Table 3. Corrosion parameters obtained from the potentiodynamic cyclic polarization curves.

Alloy	βc/ mV·dec−1	ECorr/ mV	βa/ mV·dec−1	jCorr/ µA·cm−2	EProt./ mV	EPit./ mV	RP/ Ω·cm2	RCorr/ mmpy
CrFeCuMnNi	180	−470	190	150	−220	−70	267.92	1.769
CrFeCuMnNi-3 vol.% Al2O3	170	−380	180	125	−180	−50	304.10	1.474
CrFeCuMnNi-3 vol.% Y2O3	165	−350	170	110	−135	−35	330.95	1.297
CrFeCuMnNi-3 vol.% TiO2	170	−340	180	50	−120	−30	760.25	0.589

. The values of E_Corr and j_Corr were obtained from the extrapolated Tafel lines located next to the linearized regions of the cathodic and anodic branches. The values of the protection potential, E_Prot, which is the potential after which pitting corrosion occurs, were obtained from the intersection of the reversed anodic branch with the forward anodic branch. The pitting potential, E_Pit, is the potential at which pitting corrosion starts to occur and causes a drastic increase in the current values (anodic branch side). The polarization resistance, R_P, was calculated using the following equation [32,33]:

$${\mathrm{R}}_{\mathrm{P}}=\frac{1}{{\mathrm{j}}_{\mathrm{C}\mathrm{o}\mathrm{r}\mathrm{r}}}\left(\frac{{\mathsf{\beta }}_{\mathrm{c}}.{\mathsf{\beta }}_{\mathrm{a}}}{2.3\left({\mathsf{\beta }}_{\mathrm{c}}+{\mathsf{\beta }}_{\mathrm{a}}\right)}\right)$$

(1)

Furthermore, the values of the corrosion rate, R_Corr, were obtained using the following equation [32,33]:

$${\mathrm{R}}_{\mathrm{C}\mathrm{o}\mathrm{r}\mathrm{r}}={\mathrm{j}}_{\mathrm{C}\mathrm{o}\mathrm{r}\mathrm{r}}\left(\frac{\mathrm{k}·{\mathrm{E}}_{\mathrm{W}}}{\mathrm{d}·\mathrm{A}}\right)$$

(2)

where k is a constant that specifies the unit of R_corr (k = 3272 mm (amp⁻¹ cm⁻¹ year⁻¹)), E_W is the corresponding weight of the mild steel (E_W = 28.25 g), d is the density of the mild steel (d = 7.84 gm cm⁻³), and A is the surface area of the exposed sample (A = 1 cm²).

As shown in Figure 10a, for the CrFeCuMnNi HEA, the cathodic current started to decrease with the potential until it reached the value of j_Corr. This is because of the oxygen reduction on the surface of the sample, as per the following reaction [33]:

2H₂O + O₂ + 4e⁻ = 4OH⁻

(3)

Instead, the anodic branch starts with the increase in current due to the oxidation of the surface of the alloy via its anodic reaction (via the dissolution of iron, which is the most active metal in the alloy), as follows [34]:

Fe = Fe²⁺ + 2e⁻

(4)

With a further increase in the applied potential, there is a slight decrease in the anodic current, which results from the reaction of iron with the test solution to form an oxide film, which might also be formed due to the reaction of other elements such as Ti, Cr, Cu, Mn, and Ni.

Fe + H₂O = Fe(OH)_ads + H⁺

(5)

The formation of hydroxide ions results from Reaction (5), which reacts with more Fe²⁺ to form Fe(OH)₂ on the surface of the alloy as follows:

Fe + ½ O₂ + H₂O = Fe(OH)₂

(6)

This formed ferrous hydroxide transforms to ferrous oxide (FeO) and further to magnetite (Fe₃O₄) if there is an excess of oxygen present in the solution as per the following reaction:

3Fe(OH)₂ + ½ O₂ = Fe₃O₄ +3H₂O

(7)

More applied potentials caused the current of the alloy to suddenly increase because of the breakdown of the formed oxide film and the occurrence of pitting corrosion. Here, the dissolution of the iron of the alloy in 3.5% NaCl solution into Fe²⁺ occurs under the anodic applied potential which can take place as per the following reactions [34]:

Fe + Cl⁻ = Fe(Cl⁻)_ads

(8)

Fe(OH)_ads + Fe(Cl⁻)_ads = Fe + FeOH⁺ + (Cl⁻) + 2e⁻

(9)

FeOH⁺ + H⁺ = Fe²⁺_aq + H₂O

(10)

The occurrence of pitting corrosion was indicated by a sudden increase in the current and the appearance of a hysteresis loop. This loop results from the difference between the values of the current obtained from the backward direction of the polarization curve and those obtained from the forward direction of the anodic branch. Reversing the applied potential in the reverse direction resulted in higher currents compared to the current values on the forward anodic side; the larger the size of the loop, the more severe the pitting corrosion.

The presence of Al₂O₃ within the alloy clearly had a beneficial effect on the corrosion resistance of the alloy by decreasing the values of j_Corr and R_Corr and increasing the value of R_P. This is clear from the CPP curve (Figure 10b) and the data listed in Table 3, and is due to the ability of Al₂O₃ to decrease the dissolution of the alloy surface. Moreover, the addition of Y₂O₃ within the alloy (Figure 10c) also has a powerful effect on reducing corrosion by decreasing the corrosion parameters such as j_Corr and R_Corr and increasing the value of R_P as shown in Table 3. Furthermore, TiO₂ provided the best passivation against corrosion, as shown in Figure 10d and confirmed by Table 3. The values of j_Corr and R_Corr were the lowest and the R_P values were the highest. Figure 10 also confirms the effect of Al₂O₃, Y₂O₃, and TiO₂ in reducing the severity of the pitting corrosion, where their corresponding hysteresis loops were smaller than those of the loops obtained with the alloy without the presence of any oxides. Another confirmation was given on the effect of the addition of these metallic oxides by the values of the protection and pitting potentials (E_Prot and E_Pit), which became less negative in the presence of the added oxides. The lowest negative potential values for E_Prot and E_Pit were recorded for the alloy that contains TiO₂. Thus, the polarization data indicate that the addition of metallic oxides has a positive influence on improving the corrosion resistance of the alloy, and the best performance can be listed in the following order: CrFeCuMnNi-3 vol.% TiO₂, CrFeCuMnNi-3 vol.% Y₂O₃, CrFeCuMnNi-3 vol.% Al₂O₃, and CrFeCuMnNi.

3.3.2. Electrochemical Impedance Spectroscopy (EIS) Measurements

EIS is a powerful electrochemical technique that has been successfully used to report the corrosion control of various metallic structures in harsh corrosive media [35,36,37,38]. EIS measurements were carried out to report the effect of the addition of different oxides on the corrosion passivation of the CrFeCuMnNi-fabricated high-entropy alloy in a chloride solution. Nyquist plots obtained for the (1) CrFeCuMnNi, (2) CrFeCuMnNi-3 vol.% Al₂O₃, (3) CrFeCuMnNi-3 vol.% Y₂O₃, and (4) CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs after being immersed in 3.5% NaCl solutions for 1 h before measurements are shown in Figure 11. The EIS data were also fitted to the best equivalent circuit, as displayed in Figure 12, and the values of the obtained data are listed in

Table 4. Impedance data obtained for the tested titanium alloys in the 3.5% NaCl solution.

Name of Sample	RS/Ωcm2	Q		RP1/ Ω cm2	Cdl/ F cm−2	RP2/ Ω cm2
		YQ/Fcm−2	n
CrFeCuMnNi	5.139	0.012140	0.47	29.26	1.705 × 10−5	173.9
CrFeCuMnNi-3 vol.% Al2O3	6.081	0.003579	0.54	146.5	0.001657	448.9
CrFeCuMnNi-3 vol.% Y2O3	6.172	0.002135	0.56	430.5	0.004481	1651
CrFeCuMnNi-3 vol.% TiO2	6.628	0.001826	0.60	1243	0.000380	2349

. It is worth mentioning that the employed equivalent circuit here has been used in many previous studies [39,40,41,42]. These EIS data can be defined as the solution resistance R_S, constant phase elements (CPEs) Q, polarization resistance R_P1, double layer capacitance C_dl, and another polarization resistance R_P2.

The Nyquist spectra depicted in Figure 11 show only a semicircle, and the diameter of the semicircle is the lowest for the CrFeCuMnNi HEA. The addition of Al₂O₃ (CrFeCuMnNi-3 vol.% Al₂O₃) led to increasing the diameter of the semicircle, which indicates that the presence of Al₂O₃ within the alloy increases its corrosion resistance. The addition of Y₂O₃ (CrFeCuMnNi-3 vol.% Y₂O₃) further increases the diameter but the widest diameter was obtained for the alloy that contains TiO₂ (CrFeCuMnNi-3 vol.% TiO₂). These behaviors indicate that the addition of different oxides increased the corrosion resistance of the alloy and this effect was confirmed by the values of the impedance data listed in Table 4. The values of the solution resistance, R_S, and two types of polarization resistance, R_P1 and R_P2, increased in the presence of these oxides and the highest values were recorded for the alloy containing TiO₂ (CrFeCuMnNi-3 vol.% TiO₂). The CPE under this condition has been reported previously [43], and is defined according to the following relation:

$$Z_{CPE}=\frac{\mathrm{Y}0-1}{(\mathrm{j}\mathsf{\omega })-\mathrm{n}}$$

(11)

where Y₀ is the CPE constant, ω is the angular frequency (rad S⁻¹), j² = −1 is the imaginary number, and “n” is the exponent of the CPE. It is well known that the value of “n” that accompanies the CPE (Q) determines the type of this CPE in the system under investigation. Here, the CPE can represent a resistance (Z(CPE) = R, n = 0) or capacitance (Z(CPE) = C, n = 1), or can be considered as a Warburg impedance when n = 0.5. The decrease in the value of Y_Q in the order of CrFeCuMnNi, CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% Y₂O₃, CrFeCuMnNi-3 vol.% TiO₂, and the exponent “n” (that is close to 0.5 for these experiments (0.47 < n < 0.60, as seen from Table 4, which indicates that the CPE behaves like a Warburg impedance)), indicates the passivation of the surface of the alloys due to the presence of the oxide within the alloy. Moreover, the presence of C_dl further confirms the passivation of the alloys against corrosion in the chloride solution.

To confirm the effect of the metal oxide on the corrosion behavior of the tested alloys in the chloride solution that were obtained from the Nyquist plots and polarization measurements, the Bode impedance of the interface and the degree of the phase angles were plotted, as shown in Figure 13a,b, respectively. As shown in Figure 13, the lowest recorded impedance values, |Z|, over the entire applied frequency range were obtained for the CrFeCuMnNi HEA. The addition of Al₂O₃ to the alloy led to an increase in the values of |Z| over the entire frequency range, indicating that the presence of Al₂O₃ (Figure 13, curve 2) increased the passivation of the alloy. This effect was found to increase further with the addition of Y₂O₃ (Figure 13, curve 3) and the highest |Z| values were observed for the CrFeCuMnNi- 3 vol.% TiO₂ ODS-HECs (Figure 13, curve 4) in the presence of TiO₂. It is commonly accepted [38,44,45,46] that corrosion-resistant materials provide higher |Z| values, which agrees with the obtained data and confirms that the presence of oxides within the alloy increases its passivation against corrosion in the chloride solution. The degree of the phase angle was also plotted against frequency for the tested alloys in 3.5% NaCl, as shown in Figure 13b. The minimum phase angle was recorded for the CrFeCuMnNi HEA, which had no oxides. The addition of Al₂O₃, Y₂O₃, and TiO₂ increased the maximum degree of the phase angle consecutively, indicating that the presence of these oxides increases the passivity of the alloy. This phase angle behavior confirmed that the presence of these oxides decreased the dissolution of the tested alloys and increased their passivation.

3.3.3. Chronoamperometric Current–Time (CCT) Measurements

To study the effects of the addition of different oxides on the pitting corrosion of the alloy at more anodic potentials, chronoamperometric current–time (CCP) experiments were carried out at −250 mV and −50 mV vs. Ag/AgCl. Figure 14 depicts the change in current with time curves collected at −250 mV (Ag/AgCl) for the (1) CrFeCuMnNi, (2) CrFeCuMnNi-3 vol.% Al₂O₃, (3) CrFeCuMnNi-3 vol.% Y₂O₃, and (4) CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs after their immersion in 3.5% NaCl solutions for 40 min. A value of −250 mV was chosen as the applied anodic potential from the CPP curves, which is shown in Figure 10 which reports the change in current with time. It is clear that the current of the CrFeCuMnNi HEA shows slightly steady values with time in the first few minutes of the experiment (Figure 14, curve 1), owing to the presence of an oxide or a corrosion product layer formed on its surface due to the immersion of the CrFeCuMnNi HEA for 60 min in the chloride solution before applying the anodic potential. With an increase in the duration of the experiment, the current increased almost linearly, which proves that pitting corrosion occurred for the CrFeCuMnNi HEA, the severity of which increased with the increasing duration of the experiment, which led to a continuous increase in the measured current. The addition of Al₂O₃ and CrFeCuMnNi-3 vol.% Al₂O₃ ODS-HECs (Figure 14, curve 2) decreased the absolute current with time as compared to the obtained values for the CrFeCuMnNi HEA. Furthermore, the presence of Al₂O₃ decreased the increase in the current with time, which led to a reduction in the severity of the pitting corrosion. This proves that the addition of Al₂O₃ to the alloy can passivate the alloy to some extent, which minimizes the occurrence of pitting attacks in the chloride solution. The addition of Y₂O₃ and TiO₂ to the alloy (Figure 14, curve 3 and curve 4, respectively) eliminated the occurrence of pitting corrosion, as indicated by the decrease in the current with time.

Figure 15 shows the change in current versus time which was collected at −50 mV (Ag/AgCl) for (1) CrFeCuMnNi, (2) CrFeCuMnNi-3 vol.% Al₂O₃, (3) CrFeCuMnNi-3 vol.% Y₂O₃, and (4) CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs after their immersion in 3.5% NaCl solutions for 60 min. These experiments were performed to confirm the effect of the oxide addition on the occurrence of pitting corrosion at lower negative anodic potentials. The CrFeCuMnNi HEA still shows a similar behavior to that obtained at −250 mV but with much higher values. This is because the potential value of −50 mV is much more active and aggressively dissolves the surface of the alloy in addition to causing severe pitting. The recorded current increase in the presence of Al₂O₃ (CrFeCuMnNi-3 vol.% Al₂O₃, curve 2 of Figure 15), but with highly reduced absolute values, indicates that Al₂O₃ greatly reduces the uniform corrosion, but still with a high probability of pitting corrosion occurring. The addition of Y₂O₃ within the alloy (CrFeCuMnNi-3 vol.% Y₂O₃, curve 3 of Figure 15) remarkably decreased the measured currents, proving that this oxide passivates the alloy against general corrosion. Moreover, the severity of pitting corrosion was at its minimum, as indicated by the very low increase in the current with time, particularly after 1200 s of the application of the anodic potential. The maximum effect of oxide addition was recorded for TiO₂ (CrFeCuMnNi-3 vol.% TiO₂, Figure 15, curve 4), where these currents were the lowest without any raises in their values with time. This behavior confirmed that the presence of TiO₂ provided the highest passivation against corrosion in the chloride test solution. The CCT data were in good agreement with the CPP and EIS results, indicating that the resistance to pitting corrosion was as follows: CrFeCuMnNi-3 vol.% TiO₂ > CrFeCuMnNi-3 vol.% Y₂O₃ > CrFeCuMnNi-3 vol.% Al₂O₃ > CrFeCuMnNi.

3.4. Corrosion Surface Morphology Examination

Figure 16 shows the HRSEM microstructures (surface morphology) of the corroded ODS-HEC samples captured at low (left) and high (right) magnification over the selected area. The results in Figure 16a₁,a₂ (CrFeCuMnNi HEA) reveal more pits (marked by red arrows) and oxides, indicating poor corrosion resistance. Furthermore, among the observed FCC and BCC-B2 phases, more pits were formed in the BCC-B2 area, indicating that the BCC-B2 area (rich in Fe) was more susceptible to corrosion with ${\mathrm{C}\mathrm{l}}^{-}$ than the FCC area. In other words, severe localized corrosion occurred in the unreinforced samples, owing to the large size of the grains and less structural refinement. Several authors have observed the formation of more severe corrosive pits in the BCC-B2 area compared to in the FCC area [28,47]. The HRSEM surface morphology of the CrFeCuMnNi-3 vol.% Al₂O₃ ODS-HECs sample (Figure 16b₁,b₂) exhibited a lower number of pits and oxides compared to the unreinforced CrFeCuMnNi HEA. This result confirmed that the incorporation of Al₂O₃ diminished corrosion. However, the CrFeCuMnNi-3 vol.% Al₂O₃ sample exhibited more pits and oxides compared to the CrFeCuMnNi-3 vol.% Y₂O₃ and CrFeCuMnNi-3 vol.% TiO₂ samples. The results are shown in Figure 16c₁,c₂ (CrFeCuMnNi-3 vol.% Y₂O₃), which explain that the number of pits and oxides formed had lower values compared to the CrFeCuMnNi and CrFeCuMnNi-3 vol.% Al₂O₃ samples, indicating that they were less susceptible to corrosion with ${\mathrm{C}\mathrm{l}}^{-}$. However, very few pits and almost no surface morphologies were observed in the CrFeCuMnNi-3 vol.% TiO₂ ODS-HECs sample (Figure 16d₁,d₂). This implies that the incorporation of TiO₂ enhances corrosion resistance and is less susceptible to ${\mathrm{C}\mathrm{l}}^{-}$. This was attributed to the proper embedding of TiO₂ particles over the CrFeCuMnNi matrix and greater strength due to refined grains or less grain growth compared to other samples. The HRSEM surface morphology results were consistent with or matched those of the electrochemical test, as explained in the previous section.

For further confirmation, HRSEM EDS spot analyses were carried out in different regions, which are shown in Figure 17, and the corresponding observed elemental analyses are listed in

Table 5. EDS point analysis on several regions on corroded samples for the fabricated ODS-HECs.

CrFeCuMnNi	At.%	1	2	3	4	5	6	7	8
	O	46.34	45.74	14.72	59.18	43.65	57.2	32.84	41.94
	Na	2.24	1.05	-	-	3.43	2.63	2.76	4.55
	Cl	2.93	11.19	14.32	1.15	1.53	1.5	5.86	7.72
	Cr	1.83	0.5	1.44	7.08	18.09	2.99	7.17	1.44
	Mn	9.12	7.4	18.01	4.47	4.87	1.27	10.93	5.94
	Fe	31.91	22.13	41.77	18.78	22.92	25.61	23.19	26.69
	Ni	5.08	5.52	7.18	2.71	2.89	1.07	14.93	7.99
	Cu	0.55	6.47	2.56	6.63	2.62	7.73	2.32	3.73
CrFeCuMnNi-3 vol. % Al2O3	At.%	1	2	3	4	5	6	7	8
	O	21.96	49.75	54.67	37.27	44.01	56.85	56.05	65.47
	Al	0.16	0.01	0.1	0.02	1.93	3.06	0.34	2.03
	Na	-	0.89	-	-	-	1.08	1.35	1.85
	Cl	0.28	9.92	7.1	47.31	8.68	0.22	4.85	4.2
	Cr	5.77	1.36	1.79	0.3	9.73	18.27	0.91	0.93
	Mn	7.91	5.2	2.79	13.33	5.75	9.93	3.8	3.07
	Fe	37.33	25.29	27.58	0.51	17.85	2.14	18	13.99
	Ni	13.8	4.23	4.07	0.72	5.2	1.24	6.13	4.02
	Cu	12.79	3.35	1.9	0.54	6.85	7.21	8.57	4.44
CrFeCuMnNi-3 vol. % Y2O3	At.%	1	2	3	4	5	6	7	8
	O	80.79	41.48	19.26	35.18	29.27	72.57	34.42	31.83
	Y	-	-	2.91	1.48	0.05	-	0.11	0.03
	Na	-	3.27	-	-	-	-	-	-
	Cl	1.43	0.11	0.36	0.88	3.73	3.77	15.03	15.18
	Cr	0.46	19.58	7.09	9.55	34.55	0.56	0.92	1.09
	Mn	3.51	6.5	11.3	15.35	6.95	2.76	11.74	8.43
	Fe	10.95	24.46	31.38	18.81	19.91	10.7	28.93	39.02
	Ni	2.31	2.42	14.89	9.33	2.96	6.93	5.37	3.28
	Cu	0.54	2.18	12.8	8.42	2.57	2.69	3.49	1.13
CrFeCuMnNi-3 vol. % TiO2	At.%	1	2	3	4	5	6	7	8
	O	72.55	47.06	24.57	48.61	29.42	24.79	56.87	18.56
	Ti	0.12	0.38	0.24	0.16	0.3	1.13	2.11	2.04
	Na	-	-	-	0.11	-	0.87	-	-
	Cl	2.36	12.34	1.62	1.41	0.22	1.58	2.86	0.33
	Cr	0.57	2.88	7.84	3.46	10.75	7.78	22.32	6.85
	Mn	4.59	8.06	11.27	6.2	10.79	9.27	8.43	12.45
	Fe	15.08	23.07	32.55	29.92	30.01	40.17	6.45	30.87
	Ni	4.15	5.01	10.97	2.45	11.07	5.16	0.57	15.36
	Cu	0.57	1.2	10.95	7.63	7.42	9.21	0.43	13.52

. Based on the results in Figure 17 and Table 5, it is obvious that severe corrosion was observed in the unreinforced CrFeCuMnNi sample followed by the CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% Y₂O₃, and CrFeCuMnNi-3 vol.% TiO₂ samples. More corrosive points were observed in the BCC-B2 phase area than in the FCC phase area. For example, in region 6 of CrFeCuMn HEA (Figure 17a₁), the EDS results (Table 5) exhibited O-57.2%, Na-2.63%, Cl-1.5%, Cr-2.99%, Mn-1.27%, Fe-25.61%, Ni-1.07%, and Cu-7.73%. The presence of more O and Fe atoms in this region indicated severe corrosion, which was related to Fe₂O₃. In addition, higher concentrations of Na and Cl were observed in the unreinforced CrFeCuMnNi HEA, which confirms the occurrence of severe corrosion. In region 6 of the CrFeCuMnNi-3 vol.% Al₂O₃ (Figure 17b₁), the EDS results (Table 5) exhibited O-56.85%, Al-3.06, Na-1.08, Cl-0.22, Cr-18.27, Mn-99.93, Fe-2.14%, Ni-1.24%, and Cu-7.21%. Here, the concentrations of Na and Cl atoms were lower than those of the CrFeCuMnNi sample, indicating lower corrosion. Similarly, low concentrations of Na and Cl atoms were observed in the CrFeCuMnNi-3 vol.% Y₂O₃ and CrFeCuMnNi-3 vol.% TiO₂ samples (Figure 17c₁,d₁), indicating excellent corrosion resistance. Furthermore, the presence of Al, Ti, and Y atoms in Table 5 and the corresponding EDS spectra in Figure 17b₂,c₂,d₂ clearly indicate the successful fabrication of the ODS-HECs.

4. Conclusions

In this study, CrFeCuMnNi, CrFeCuMnNi-3 vol.% Al₂O₃, CrFeCuMnNi-3 vol.% TiO₂, and CrFeCuMnNi-3 vol.% Y₂O₃ of oxide-dispersed high-entropy composites were successfully developed via hot forging. A homogeneous supersaturated solid solution of the incorporated elements was successfully achieved after 50 h of MA, which was evidenced by the SEM powder surface morphology. The developed materials exhibited a major FCC phase and a minor BCC-B2 phase after hot forging, which was confirmed via X-ray diffraction and SEM microstructures. Fewer grain growths and refined structures were obtained for the CrFeCuMnNi-3 vol.% TiO₂ sample due to uniform dispersion and easy adherence with the matrix. The corrosion performance of the developed materials was evaluated using electrochemical tests in 3.5% NaCl solution. The results demonstrated that the incorporated oxide particles enhanced corrosion resistance. Among the four novel materials developed, the CrFeCuMnNi-3 vol.% TiO₂ sample exhibited the highest corrosion resistance compared to the other samples. The optimal corrosion properties of the CrFeCuMnNi-3 vol.% TiO₂ sample produced anticorrosion performance of a low value of corrosion current (j_Corr) of 50 μA/cm², corrosion potential (E_Corr) of −340 mV, and corrosion rate (R_Corr) of 0.589 mmpy. These results suggest that the developed materials are suitable for industrial application.