Development of FeCrNi medium entropy alloys with excellent mechanical properties and corrosion resistance

ABSTRACT

The development of low-cost high-entropy alloys is promising to address the demands of expanding industrial applications. In the present research, Co-free FeNiCr medium-entropy alloys were prepared by arc melting and post-processing treatments. The influence of Fe content on the microstructure, mechanical properties and corrosion resistance of these alloys were clarified. The results show that the FeNiCr medium entropy alloys exhibit a single face-centered cubic structure. These alloys consist of a multitude of fine grains and a high density of annealing twining boundary distributed at the interface of the larger grains. The Fe40Ni30Cr30 exhibits both high strength and excellent elongation. The excellent mechanical properties of FeNiCr alloys are mainly attributed to the solid solution strengthening, grain boundary strengthening, twinning induced plasticity (TWIP) effect, dimples toughening, and micropores toughening. The increase in Fe content has been observed to impede the segregation of Cr, dislocation pinning, and the nucleation of grains at the grain boundaries, leading to an increase in elongation and a reduction in strength. In addition, the electrochemical experimental results show that a stable passivation film is formed on the surface of Fe40Ni30Cr30 alloy during the corrosion process, which reduces the corrosion rate and improves the corrosion resistance of the alloy.

Keywords:
Medium entropy alloys; Microstructure; Mechanical properties; Corrosion resistance

1. INTRODUCTION

The development of cost-effective HEAs with comprehensive properties is critical to promote industrial applications of HEAs in extreme environments. In the past few decades, some cost-effective Fe-rich high-entropy alloys have been reported to exhibit excellent properties and have attracted widespread attention [¹], such as FeNiCrMo [²], FeMnCoCr [³], FeNiCrAl [⁴], FeCrNiMnCu [⁵], etc. Since the entropy values of most alloys are not at a high entropy level, they are generally classified as medium entropy alloys according to the classification of YEH et al. [⁶]. Nevertheless, the entropy of the Fe-rich alloys is indeed higher than traditional steels due to their multicomponent characteristic.

According to literature research, the methods reported to improve the mechanical properties of high entropy alloys mainly include solid solution strengthening, dislocation strengthening deduced from pre-existing dislocations, grain boundary strengthening, precipitation strengthening, twin boundary strengthening, and phase-transformation induced strengthening, etc [⁷, ⁸, ⁹]. Most medium-entropy or high-entropy alloys with FCC structure are usually prone to TWIP or TRIP effect during deformation due to their low stacking fault energy [¹⁰]. LI et al. [³] reported a Fe₅₀Mn₃₀Co₁₀Cr₁₀ medium-entropy alloy that exhibited excellent strength and plasticity. In order to break the trade-off between strength and ductility, it is usually desirable to obtain a dispersed strengthening phase that is coherent with the matrix in high entropy alloys [¹¹]. Adding elements such as Mo, Al, and Ti to cost-effective Fe-rich medium-entropy alloys is a common method to improve the strength of these alloys. Yu YIN et al. [¹²] reported that the addition of Mo led to the formation of σ phase and μ phase in FeNiCrMo medium entropy alloys. Since σ phase and μ phase are typical brittle phases, although the compressive strength of Fe₄₅Ni₂₅Cr₅Mo₂₅ alloy reaches 1661 Mpa, the ductility of the alloy drops significantly to 17%. Adding Ti and C to FeNiCrMo alloy results in the in-situ precipitation of carbide particles during aging, which significantly improves the strength, hardness and wear resistance of the alloy. However, the plasticity of the alloy decreases significantly [¹³]. In addition, heterogeneous structure is also one of the effective ways to improve the strength and toughness of alloys [¹⁴].

Currently, the commonly used methods for preparing cost-effective medium-entropy alloys include powder metallurgy, mechanical alloying, and arc melting, etc [¹⁵, ¹⁶, ¹⁷]. Homogenization annealing, rolling and stress relief annealing are effective methods to obtain alloys with uniform composition and fine grain size [¹⁷]. Researches have shown that modulating Fe content in FeCoNiCr based MEAs has the effect of enhancing tensile properties [¹⁸], fracture toughness [¹⁹], oxidation resistance behavior [²⁰], and compressive properties [²¹]. ZENG et al. [²²] reported that the tensile strength and elongation of the AlCoCr_0.5FexNi_2.5 (x = 0.5, 1.5, 2.5, and 3.5) alloy increased by 23% and 133%, respectively, when the proportion of Fe content was increased from 0.5 to 1.5. JEONG and PARK [²³] reported that further increase Fe content in the Fex(CoCrMnNi)₁₀₀₋x MEAs from Fe65 to Fe70 significantly strengthened the alloy system, resulting in reduced ductility. POVOLYAEVA et al. [¹⁹] reported similar results in Fe_x(CoNi)_100−xCr_9.5C_0.5 HEAs. FAN et al. [²⁴] suggested that adding too much Fe (≥ 47.5 at. %) into AlCoCrFeNi HEA coating will deteriorate the mechanical and corrosion properties. Further, the high-density grain boundary and the Fe-rich passive layer showed a negative impact on the corrosion behavior of ultra-fine grained CoNiFeCrMn HEAs [²⁵]. Up to now, there is no consensus in the literature regarding the effect of Fe on the mechanical properties and corrosion resistance of cost-effective medium-entropy alloys. Most of the research results are affected by the precipitation phases, such as σ phase, μ phase, carbides, BCC martensite, etc. Moreover, the corrosion resistance of cost-effective FeCrNi based MEAs has been largely overlooked in the existing literatures.

In order to further clarify the effect of Fe element on the mechanical properties and corrosion resistance of cost-effective medium-entropy alloys and eliminate the influencing factors of precipitation phase, three distinct FeNiCr MEAs with a face-centered cubic (FCC) single-phase structure were successfully prepared, namely Fe_33.3Ni_33.3Cr_33.3, Fe₄₀Ni₃₀Cr₃₀, and Fe₅₀Ni₂₅Cr₂₅. The effect of Fe content on the microstructure, mechanical properties, and corrosion resistance of FeNiCr MEAs was investigated in detail.

2. EXPERIMENTAL PROCEDURES

2.1. Materials and preparation of FeNiCr MEAs

The high-purity Fe, Ni and Cr bulk metals are utilized as the raw materials for preparation of alloy ingots through vacuum arc melting (99.99%, Beijing Dream Material Technology Co., Ltd., China). The nominal composition of alloys is delineated in Table 1. Each specimen was remelted a minimum of six times during the melting process with electromagnetic stirring employed throughout. A button ingot of approximately 60 g of MEAs was produced. Subsequently, the ingots were subjected to a series of post-processing treatments as follow. The ingots were vacuum sealed in a high-purity argon environment and placed in a muffle furnace at 1200 °C for 12 h for homogenization annealing, followed by water quenching treatment. Then, hot rolling treatment was performed at 800 °C to refine the grain of homogenized alloys with a deformation of 60%. Following hot rolling, the samples were subjected to cold rolling at room temperature with a deformation of 10%. Subsequently, the samples were placed in a muffle furnace for stress relief annealing at 700 °C for 30 minutes.

Table 1
Nominal composition (at.%) of FeNiCr MEAs.

MATERIALS	Fe	Ni	Cr
Fe33.3Ni33.3Cr33.3	33.33	33.33	33.33
Fe40Ni30Cr30	40	30	30
Fe50Ni25Cr25	50	25	25

2.2. Microstructure characterization

The phase structure of FeNiCr MEAs was identified by using an X-ray diffractometer with a Cu Kα radiation (D/max 2550VB+, Rigaku Corporation, Japan). The voltage and current were set to 40 kV and 200 mA, respectively. The diffraction angle (θ) was set within the range of 10° to 90°, with a scanning step of 0.02°.

The microstructure of FeNiCr MEAs was observed by using a high-resolution field emission scanning electron microscope (SEM, Quanta 250 FEG, FEI, USA) with an electron backscattering diffraction (EBSD) system. Before SEM testing, the surface of these alloys was polished, then subjected to grain boundary etching using an etching solution comprising 1 g FeCl₃, 20 ml HCl, and 100 ml H₂O. The specimens for EBSD analysis were prepared by mechanical polishing, followed by electro-polishing using a mixture solution of 10 vol % perchloric acid and methanol solution with the voltage of 20 V. The transmission electron microscopy (TEM, Tecnai F20, FEI, USA) was used to further investigate the microstructure at an operating voltage of 200 kV.

2.3. Tensile testing

Flat dog-bone shape specimens with gauge dimensions of 26 mm × 10 mm × 2 mm were cut from plates of each sample and were polished on both sides to mirror finish. The FeNiCr MEAs samples were subjected to tensile testing at room temperature with a test strain rate of 1×10₋³ s₋¹ by an Instron universal testing machine. The representative data for each condition were obtained by averaging three values of the test results.

2.4. Corrosion resistance testing

The electrochemical tests were conducted on a multichannel electrochemical workstation with a three-electrode cell (Multi Autolab M204, Aptar, Switzerland). The core parameters of the workstation are shown in Table 2.

Table 2
Core parameters of electrochemical workstation.

PARAMETERS	VALUE
Potential accuracy	±0.2%
Current accuracy	±0.2%
Current	10 mA – 10 nA
Potentiostatic	±10V
Frequency	10 μHz – 1 MHz

The tested sample with an exposed area of 10 × 10 mm was polished and employed as the working electrode, a platinum electrode was utilized as the counter electrode, and a saturated calomel electrode (SCE) was employed as the reference electrode. All electrochemical tests were conducted in a 3.5 wt.% NaCl solution at room temperature. The working electrode was immersed in the electrolyte at open circuit potential (OCP) for 30 minutes to allow for the attainment of an electrochemical steady-state surface. The signals of electrochemical impedance spectroscopy (EIS) were recorded over a scan frequency ranging from 10 mHz to 100 kHz with an amplitude of 10 mV. The polarization curves were recorded at a scan rate of 1.0 mV/s over the potential range of −0.5 V_SCE to 1.5 V_SCE. At least three specimens were used for each test to ensure the reliability. The microstructure of corrosion surface of FeNiCr MEAs was observed using a field emission scanning electron microscope (Tescan Mira4, Tescan, Brno, Czech Republic) equipped with a backscattered electron diffractometer.

3. RESULTS AND DISCUSSION

3.1. Microstructure

Figure 1 presents the XRD patterns of FeNiCr MEAs. It can be observed that all the FeNiCr MEAs exhibit FCC single-phase structure. No second phase is observed, indicating that Fe, Ni, and Cr elements form a simple solid solution due to the high mixing entropy. Furthermore, the increase in Fe content does not alter the phase structure of the solid solution.

Figure 1
X-ray diffraction patterns of FeNiCr MEAs.

The morphologies of FeNiCr MEAs in the perpendicular and parallel to the rolling direction after post-processing treatments are shown in Figure 2. Following the erosion of the FeCl₃+ HCl corrosive solution, the grain boundaries of the large grains in the alloy are clearly evident (as indicated by the white arrows). The large grains observed in the Fe_33.3Ni_33.3Cr_33.3 show a markedly large grain size, measuring over 200 µm. As the Fe content increases, the number of interfaces with large grains decreases, and the presence of interfaces is even difficult to observe in Fe₅₀Ni₂₅Cr₂₅ (Figure 2(c)). Furthermore, alternating corrosion steps were identified in the Fe_33.3Ni_33.3Cr_33.3 (Figure 2(a)), which may represent a complex structure formed in the homogenized plate-like organization following large deformation rolling. Table 3 depicts the outcomes of the elemental analysis of the three alloys. The EDS results suggest that the elemental compositions of the MEAs are comparable to their nominal compositions.

Figure 2
Microstructure of FeNiCr MEAs perpendicular (a~c) and parallel (d~f) to the rolling direction, (a, d) Fe_33.3Ni_33.3Cr_33.3, (b, e) Fe₄₀Ni₃₀Cr₃₀, (c, f) Fe₅₀Ni₂₅Cr₂₅.

Table 3
Composition (at.%) of points in Figure 2.

MATERIALS	POSITION	Fe	Ni	Cr
Fe33.3Ni33.3Cr33.3	1	33.34 ± 3.57	30.46 ± 4.05	36.20 ± 2.97
Fe40Ni30Cr30	2	40.94 ± 3.37	26.96 ± 4.58	32.10 ± 3.24
Fe50Ni25Cr25	3	50.87 ± 3.00	22.42 ± 5.21	26.70 ± 3.56

Figures 2(d~f) provide insight into the microstructure of these alloys, illustrating its characteristics parallel to the rolling direction. As illustrated in Figures 2(d) and (e), the interface of the large grains is constituted by a multitude of fine grains, which extend from the interface to the interior of the large grains. In the Fe₄₀Ni₃₀Cr₃₀, fine grains are observed exclusively at the interface of the large grains (Figure 2(e)).

The morphology of fine grains at the interface of large grains in the Fe₄₀Ni₃₀Cr₃₀ MEA was observed by EBSD, as illustrated in Figure 3. In Figure 3(a), some large grains with distinct colors are evident, and a considerable number of fine recrystallized grains are dispersed at the grain boundaries and polycrystalline intersections. The grain orientations are randomly distributed, exhibiting no discernible texture. The grains in the FeNiCr MEAs can be classified into three types. The first comprises ultra-coarse equiaxial grains. These are mainly large grains formed by homogenization treatment and non-obviously oriented grains formed by rolling treatment. The second comprises a small number of deformed grains formed at the grain boundaries of ultra-coarse equiaxial grains. These are supposed to be annealed twins and secondary crystallization grains formed by the rolling-annealing process [²⁶]. These grains with annealing twining boundary can be observed in Figures 3(c) and (d), where they are situated at the interface between ultra-coarse equiaxial grains. The annealing twining boundary exhibits a significant concentration of 60° orientation of Σ3<111>/60°, accounting for approximately 21% of the total content. The formation of annealing twins may be related to dislocations movement and the high-temperature annealing process. The high-temperature annealing process caused a decrease in the density of dislocations and an increase in the annealing twins [²⁷]. In addition, ultrafine recrystallized grains were found at the interface of the large grains with a grain size of about 1 µm. The dislocations and residual stresses stored at the grain boundaries provide the driving force for nucleation of recrystallized grains, which subsequently grow during the annealing process [²⁸].

Figure 3
Microstructure of Fe₄₀Ni₃₀Cr₃₀ MEA, (a) IPF image, (b) distribution diagram of grain size, (c) distribution of twining boundaries, and (d) distribution map of misorientation.

The TEM morphology of Fe₄₀Ni₃₀Cr₃₀ MEA is shown in Figure 4. It can be observed from Figure 4a that the alloy contains a large number of micron-sized grains. As illustrated in Figure 4(b), the alloy exhibits a considerable number of dislocations. During the rolling process, the generation and movement of dislocations are impeded by grain boundaries, leading to their accumulation at the grain boundaries. The annealing process causes the accumulated dislocations to evolve into twins and simultaneously encourages their growth to form fine grains. The local chemical heterogeneity results in a larger atomic size difference, leading to significant lattice distortion and dislocation density in HEA [²⁹]. It can also be found from Figures 4c and 4d that the alloy contains more FCC structure grains with different orientations.

Figure 4
Microstructure of Fe₄₀Ni₃₀Cr₃₀, (a) bright field TEM image of fine grain region, (b) high magnification TEM image, (c) TEM image at the interface between two fine grains, and (d) diffraction pattern of interface.

3.2. Mechanical properties

3.2.1. Tensile properties

Figure 5 presents the stress-strain curves of FeNiCr MEAs. The yield strength,tensile strength and ductility for these alloys is provided in Table 4. The Fe_33.3Ni_33.3Cr_33.3 exhibits the highest tensile strength (878.40 MPa), yet its plasticity is relatively low at 18.04%. As the proportion of Fe in the alloy increases, both yield strength and tensile strength of MEA decrease while its plasticity increases. The yield strength of Fe₄₀Ni₃₀Cr₃₀ is 628.78 MPa, the tensile strength is 776.7 MPa, and the elongation is 26.88%. In comparison to Fe_33.3Ni_33.3Cr_33.3, Fe₄₀Ni₃₀Cr₃₀ exhibits a 15% reduction in yield strength and an 11% decline in tensile strength, accompanied by a 49% augmentation in elongation. Despite a 35% reduction in elongation compared to Fe₅₀Ni₂₅Cr₂₅, the yield strength and tensile strength of Fe₄₀Ni₃₀Cr₃₀ exhibited a 48% and 19% increase, respectively. In comparison to the three alloys, the Fe₄₀Ni₃₀Cr₃₀ demonstrated an optimal equilibrium between strength and plasticity. Furthermore, the developed FeNiCr MEAs exhibit superior tensile properties compared to the majority of other MEAs/HEAs, as illustrated in Figure 6 [¹⁵, ¹⁹, ²², ²³, ³⁰, ³¹, ³², ³³, ³⁴, ³⁵].

Figure 5
Stress strain curves of FeNiCr MEAs.

Table 4
Mechanical properties of FeNiCr MEAs.

MATERIALS	YIELD STRENGTH (MPa)	TENSILE STRENGTH(MPa)	DUCTILITY (%)
Fe33.3Ni33.3Cr33.3	739.76	878.40	18.04
Fe40Ni30Cr30	628.78	776.70	26.88
Fe50Ni25Cr25	424.02	651.31	41.73

Figure 6
Comparison of tensile properties between FeNiCr MEAs and other medium/high entropy alloys [¹⁵, ¹⁹, ²², ²³, ³⁰–³⁵].

3.2.2. Discussion on strengthening mechanism

Figure 7 shows the EBSD analysis results of Fe₄₀Ni₃₀Cr₃₀ MEA in the area proximate to the fracture site. As illustrated in Figures 7(a) and (c), a considerable number of elongated twining grains and slip bands have emerged in the Fe₄₀Ni₃₀Cr₃₀ MEA following tensile deformation, indicating that the initial organization of annealed twins has been significantly disrupted. Figure 7(b) demonstrates that there is no discernible difference in grain size before and after the fracture of the specimen in the undeformed region. The presence of fine grains and grain boundaries may effectively resist deformation, enhancing the strength. The annealed twins of deformed Fe₄₀Ni₃₀Cr₃₀ are significantly diminished, suggesting that the annealed twins as effective coordinators in the tensile deformation process. During tensile deformation, annealed twins underwent a gradual transformation into deformed twins, relieving stress concentration, promoting uniform plastic deformation, and enhancing the plasticity [³⁶]. In conclusion, the excellent balance of strength and toughness of Fe₄₀Ni₃₀Cr₃₀ is mainly attributed to the combined effects of solid solution strengthening, grain boundary strengthening, and TWIP effect.

Figure 7
EBSD analysis results of Fe₄₀Ni₃₀Cr₃₀ MEA after tensile fracture: (a) IPF image, (b) diagram of grain size distribution, (c) distribution of twining boundaries, and (d) distribution map of misorientation.

Figure 8 shows the fracture morphologies of FeNiCr MEAs after tensile test. The predominant fracture mode observed in FeNiCr MEAs is shear fracture. The fracture features of Fe_33.3Ni_33.3Cr_33.3 are predominantly dimples, which encompass a limited number of microporous features and cleavage fracture of large grains (Figures 8(a) and (d)). The latter can be attributed to the fracture of brittle grains. As the Fe content increased, the density and number of dimples increased substantially, and the cleavage fracture gradually disappeared. In the Fe₄₀Ni₃₀Cr₃₀, the area of the cleavage fracture exhibits a marked decrease, concomitant with the observation of a low density of micropores, as illustrated in Figures 8(b) and (e). High density of dimples and fine micropores are characteristic of the ductile features were observed in the fracture surface of Fe₅₀Ni₂₅Cr₂₅, contributing to the high ductility. In comparison to the Fe₄₀Ni₃₀Cr₃₀, the microporous size on the Fe₅₀Ni₂₅Cr₂₅ increases, including some large microporous with a size of about 5 µm. In the fracture process, the micropores initially manifested at the bottom of the dimples. The tensile stress then caused the dislodgement or fracturing of the dimples along a direction perpendicular to the tensile stress, resulting in the nucleation of micropores. The micropores commenced a process of diffusion and growth with the onset of plastic deformation, leading to the formation of larger micropores. The thin walls of these large micropores finally necked and destructed by shear stress [³⁷]. An increase in the Fe content, an increase in the density and size of the micropores, improving the shear resistance and elongation of MEAs.

Figure 8
The fracture morphologies of FeNiCr MEAs, (a, d) Fe_33.3Ni_33.3Cr_33.3, (b, e) Fe₄₀Ni₃₀Cr₃₀, (c, f) Fe₅₀Ni₂₅Cr₂₅.

The strengthening mechanisms of FeNiCr MEAs can be broadly classified into four categories: solid solution strengthening, grain boundary strengthening, twinning-induced plasticity and transformation induced plasticity. The above strengthening and toughening mechanisms have been characterized and discussed in detail in our previous work on the deformation behavior of powder metallurgy and 3D printed Fe_33.3Ni_33.3Cr_33.3 alloys [³⁸, ³⁹]. The solid solution strengthening mechanism is one of the primary recognized strengthening mechanisms for MEAs [⁴⁰]. The generation and movement of dislocations are impeded by the severe lattice distortion effect, thereby enhancing the strength of medium-entropy alloys. This phenomenon has been experimentally validated by Ibrahim Ondicho and colleagues [⁴¹]. Their findings indicate that the incorporation of Fe into CoCrMnNiFe high-entropy alloys mitigates the extent of lattice distortion and the solid solution strengthening effect. Thus, one of the primary factors contributing to the relatively low strength of FeNiCr MEAs with a higher Fe content is the diminished extent of lattice distortion. Further, the results of morphological structure observation indicate that grain boundary strengthening is one of the primary mechanism responsible for the high strength of FeNiCr MEAs. The fine grain boundaries impede the movement of dislocations, making it more challenging for them to traverse the material. Consequently, the refinement of grains and the increase in the number of grain boundaries collectively enhance the strength of the material. As a result of rolling and subsequent annealing in the FeNiCr MEAs, a considerable number of annealing twins and twining boundaries were introduced, thereby further impeding the movement of dislocations and improving the strength. This illustrates that strengthening mechanism of FeNiCr MEAs encompasses solid solution strengthening, grain boundary strengthening, and TWIP effect.

3.3. Corrosion resistance

3.3.1. Potentiodynamic polarization curves

Figure 9 illustrates the potentiodynamic polarization curves of the FeNiCr MEAs in a 3.5 wt.% NaCl solution. The Fe₅₀Ni₂₅Cr₂₅ MEA exhibits two distinct passivation stages, with the potential at which passivation occurs being −0.23 V and 1.24 V, respectively [⁴²]. The potentiodynamic polarization curve of the Fe₅₀Ni₂₅Cr₂₅ MEA exhibit numerous minor current fluctuations, indicating the generation of sub-stable pits and low stability of the passivation film on the surface of the alloy. In contrast, the Fe_33.3Ni_33.3Cr_33.3 and Fe₄₀Ni₃₀Cr₃₀ MEAs do not demonstrate any discernible passivation behavior. The Fe₄₀Ni₃₀Cr₃₀ MEA displays a paucity of current fluctuations. The electrochemical parameters, including the corrosion potential (E_corr) and corrosion current density (I_corr), were derived from Figure 9, as presented in Table 5. The Zview program was employed to calculate the anodic Tafel slope (β_a) and the cathodic Tafel slope (β_c). The Stern-Geary equation (Formula 1) is employed for the calculation of the polarization resistance (Rp).

Figure 9
Potentiodynamic polarization curves of FeNiCr MEAs in 3.5 wt.% NaCl solution.

Table 5
Electrochemical parameters calculated from the potentiodynamic polarization curves.

	ECORR (mV)	ICORR (A·cm−2)	βa (V·dec−1)	βc(V·dec−1)	Rp(Ω·cm2)
Fe33.3Ni33.3Cr33.3	−194	6.5 × 10−8	0.33	−0.11	551120
Fe40Ni30Cr30	−223	5.5 × 10−8	0.32	−0.09	554564
Fe50Ni25Cr25	−240	8.7 × 10−8	0.44	−0.1	400021

(1) $$R_p=\frac{dE}{di}=\frac{1}{i_{corr}}\frac{{\beta }_a{\beta }_c}{2.303({\beta }_a+{\beta }_c)}$$

Where i_corr is the corrosion current density. A high corrosion potential and low corrosion current density indicate that the alloy has high chemical stability and low corrosion tendency. As evidenced in Table 5, the Fe₅₀Ni₂₅Cr₂₅ MEA exhibits the lowest corrosion potential, the highest corrosion current density, and the smallest polarization resistance, indicating that it exhibits the poorest corrosion resistance. The corrosion potentials of the Fe_33.3Ni_33.3Cr_33.3 and Fe₄₀Ni₃₀Cr₃₀ MEAs are relatively similar. Nevertheless, the Fe₄₀Ni₃₀Cr₃₀ MEA exhibits a lower corrosion current density and a higher polarization resistance. This indicates that the Fe₄₀Ni₃₀Cr₃₀ MEA exhibits low corrosion rate and excellent corrosion resistance.

3.3.2. Electrochemical impedance spectroscopy

Figure 10 illustrates the electrochemical impedance spectroscopy (EIS) of the FeNiCr MEAs under OCP conditions. The Nyquist plots of all three MEAs in the 3.5 wt.% NaCl solution are incomplete semicircles (Figure 10(b)), which suggest that the corrosion process of these alloys is controlled by the charge transfer process. The radius of the semicircle is proportional to the charge transfer resistance (R_film) of the electrode. The slightly larger radius of the Nyquist plots for the Fe₄₀Ni₃₀Cr₃₀ MEA indicates a higher R_film, and enhanced corrosion resistance against electrochemical dissolution. Figure 10(c) depicts the logZ-logf curves in the Bode plots of the three HEAs, which exhibit a linear relationship with a negative slope over a broad frequency range. This suggests that the passivation film formed on these alloys possesses pseudocapacitive properties and that a chemical energy storage reaction occurs in the corrosion solution. Furthermore, the Fe₄₀Ni₃₀Cr₃₀ exhibits the highest impedance value. In the mid-frequency region of the Bode plot, the angle θ is approximately −80°, which suggests the formation of a stable passivation film in a 3.5 wt.% NaCl solution [⁴³]. As illustrated in Figure 10(d), the phase-logf curves of FeNiCr MEAs exhibit a nearly horizontal straight line across a wide frequency range. This observation suggests that the passivation film formed on the surface of these alloys possesses capacitance-like properties and exhibits a high capacitive resistance. The R_s(Q_film(R_film(Q_dlR_ct))) equivalent circuit, as proposed in the literature, was employed to fit the EIS data of this study using Zsimpwin software [⁴⁴]. The equivalent circuit diagram is shown in Figure 10(a), and the fitted parameters are listed in Table 6. R_s, R_film, and R_ct represent the solution resistance, passivation film resistance and charge transfer resistance, respectively. Additionally, Q_film and Q_dl denote the equivalent capacitance of the passivation film and electric double layer, respectively.

Figure 10
EIS results of FeNiCr MEAs in 3.5 wt% NaCl solution under open circuit potential condition, (a) the equivalent circuit used for fitting EIS experimental data, (b) Nyquist plots and (c, d) Bode plots.

Table 6
The fitting results of the EIS spectra.

	Rs (Ohm·cm2)	FILM		DOUBLE LAYER		Rfilm(Ohm·cm2)	Rct(Ohm·cm2)	Qfilm (Fcm−2)	Qdl (Fcm−2)
		Y0(Ω−1·cm−2·s−n)	n	Y0(Ω−1·cm−2·s−n)	n
Fe33.3Ni33.3Cr33.3	1.5	6.7 × 10−6	0.84	2.6 × 10−6	0.8	3.6 × 105	2.5 × 106	7.9 × 10−6	1.6 × 10−7
Fe40Ni30Cr30	1.4	2.9 × 10−6	0.86	2.1 × 10−6	0.6	5.6 × 105	3.7 × 106	3.1 × 10−6	4.4 × 10−10
Fe50Ni25Cr25	2.0	7.8 × 10−6	0.88	1.9 × 10−6	0.53	3 × 105	2.2 × 106	8.8 × 10−6	3 × 10−11

As the passivation films on the surfaces of the samples are not homogeneous, their capacitance is represented by the equivalent capacitance (Q) [⁴⁵]. The equivalent capacitance Q related to CPE can be calculated using Brug’s formula [⁴⁶].

(2) $$Q=Y_0^{1/n}{R}_s^{(1-n)/n}$$

The parameter Q consists of two components: the constant Y₀ and the dimensionless exponent n. The value of n represents the deviation from purely capacitive behavior, with a closer value of n to 1 indicating a closer alignment with the ideal capacitance [⁴⁷]. The value of n is less than 1 for all MEAs, indicating that the electrochemical behavior of the passivated film deviates from purely capacitive behavior [⁴³]. By selecting R_s and R_film, the equivalent capacitances Q_dl and Q_film of the double layer and the passivation film can be calculated according to the Brug formula, as shown in Table 6 [⁴⁸].

In general, equivalent capacitance represents the ability to carry charge. The smaller the capacitance value, the fewer anions the double-layer structure can bond or absorb, thereby reducing the damage to the passivation film. The thickness of the passivation film is inversely proportional to the capacitance value [⁴⁹]. The thickness of the passivation film can be estimated according to formula (3) [⁴⁹].

(3) $$d_{eff}=\frac{\epsilon {\epsilon }_0}{Q_{film}}$$

Where ε = 12 is the relative dielectric constant of the passivation film [⁵⁰]. ε₀ is the dielectric constant of vacuum (8.8542 × 10⁻¹⁴ F/cm). The calculation results show that the thickness of the passivation films of Fe_33.3Ni_33.3Cr_33.3, Fe₄₀Ni₃₀Cr₃₀ and Fe₅₀Ni₂₅Cr₂₅ MEA are 1.3 nm, 3.4 nm and 1.2 nm respectively.

As evidenced in Table 6, the R_s values of the MEAs are relatively similar and do not impact the overall assessment of performance. The Fe₄₀Ni₃₀Cr₃₀ exhibits the smallest Y₀ value (2.9 × 10⁻⁶ Ω⁻¹·cm⁻²·s⁻ⁿ), indicating that the passivation film on its surface is of greater thickness and density. The passivation film exhibits a high charge transfer resistance and affects the charge transfer resistance value [⁵¹]. As the R_film value increases, the ability of the charge to traverse through the passivation film is concomitantly constrained [⁵²]. The R_film value of the Fe₄₀Ni₃₀Cr₃₀ MEA is more than 1.5 times that of the Fe_33.3Ni_33.3Cr_33.3 MEA and nearly 2 times that of the Fe₅₀Ni₂₅Cr₂₅ MEA. And the Fe₄₀Ni₃₀Cr₃₀ also displays a higher R_ct value, indicating more stable passivation film, less sensitive to the anions in the solution, and a lower rate of chemical reactions.

Figure 11 shows the microstructure of the passive film on the surface of three medium-entropy alloys after corrosion in 3.5 wt.% NaCl solution. It can be seen from Figures 11(a) and (d) that local corrosion areas appeared on the surface of Fe_33.3Ni_33.3Cr_33.3 MEA. Corrosion preferentially initiates at grain boundaries. The passivation film formed on the surface of the alloy after corrosion is uneven and has tiny holes on its surface. It can be seen from Figure 11(b) that the corrosion area on the surface of the Fe₄₀Ni₃₀Cr₃₀ medium entropy alloy increases. However, as shown in Figure 11(e), the passive film formed on the alloy surface is dense and smooth. There are no obvious defects on the surface of the passive film. Granular corrosion products can also be observed on the surface of the passivation film. On the contrary, it can be found from Figures 11(c) and (f) that a large number of corrosion pits appeared on the surface of the corroded Fe₅₀Ni₂₅Cr₂₅ alloy. The passive film on the alloy surface has a large number of holes. In addition, a large amount of corrosion products were also found in the holes. As shown in Figure 9, the anodic polarization curve of the F_e50Ni₂₅Cr₂₅ alloy during the corrosion process showed a lot of fluctuations. This indicates that the passivation film formed on the surface of the alloy during the corrosion process is not dense. After the passive film was broken through, repassivation occurred on the alloy surface. This is consistent with the morphological observation results shown in Figure 11. Table 7 gives the EDS analysis results for the marked positions in Figure 11. The results showed that the passivation films on the surfaces of the three medium-entropy alloys were all composed of mixed oxides and chlorides of Fe/Ni/Cr. The proportion of Ni and Cr elements in the passivation film on the surface of Fe₄₀Ni₃₀Cr₃₀ medium entropy alloy is significantly higher than that of the other two alloys. In summary, the passivation film formed on the surface of Fe₄₀Ni₃₀Cr₃₀ MEA is dense and smooth. The alloy exhibits excellent corrosion resistance.

Figure 11
Surface morphologies of the passivation film on the surface of FeNiCr MEAs after corrosion in 3.5 wt.%NaCl solution. (a, d) Fe_33.3Ni_33.3Cr_33.3, (b, e) Fe₄₀Ni₃₀Cr₃₀, (c, f) Fe₅₀Ni₂₅Cr₂₅.

Table 7
EDS analysis results for the marked positions in Figure 11 (at.%).

POSITION	Fe	Ni	Cr	Na	Cl	O
P1	32.57	33.56	28.96	0.00	0.01	4.89
P2	31.58	31.23	33.27	0.20	0.03	3.69
P3	36.18	26.59	28.32	5.09	1.19	2.64
P4	38.08	28.10	29.72	1.04	0.23	2.82
P5	46.88	24.31	24.12	0.72	1.16	2.81
P6	38.76	16.81	22.12	9.39	1.00	11.92

4. CONCLUSIONS

The FeNiCr MEAs with single FCC phase were prepared through casting method and post-processing treatments. The effect of Fe on the microstructure, mechanical properties and corrosion resistance of FeNiCr medium entropy alloys was systematically studied. The main conclusions are as follows.

1. All the FeNiCr MEAs exhibit a single-phase FCC structure. The increase in Fe content does not change the phase structure of the FCC solid solution. Fe₄₀Ni₃₀Cr₃₀ MEA exhibits a heterogeneous structure consisting of large-sized grains and small-sized grains. In addition, there are many annealing twins inside the alloy.

2. With the increase of Fe content, the strength of FeNiCr MEAs gradually decreases, while the elongation of these alloys gradually increases. Fe₄₀Ni₃₀Cr₃₀ exhibits both high strength and elongation, with a yield strength, tensile strength, and elongation of 628.78 MPa, 776.70 MPa, and 26.88%, respectively. The high strength of FeNiCr MEAs can be attributed to solid solution strengthening, grain boundary strengthening, and the TWIP effect.

3. Low corrosion current density and high polarization resistance indicate that the Fe₄₀Ni₃₀Cr₃₀ MEA has a low corrosion rate. The passive film formed on the surface of the alloy during the corrosion process is less sensitive to anions, resulting in excellent corrosion resistance.

5. ACKNOWLEDGMENTS

The authors acknowledge the National Natural Science Foundation of China (No. 52101208), Natural Science Foundation of Hunan Province of China (2022JJ50175), Scientific Research Foundation of Hunan Provincial Education Department (21A0465), and Hunan Province Graduate Research Innovation Project (CX20240993).

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Publication Dates

• Publication in this collection
  16 May 2025
• Date of issue
  2025

History

• Received
  07 Feb 2025
• Accepted
  11 Apr 2025