Unveiling the tribological and corrosion characteristics of Ti20-Al16-V16-Fe16-Ni16-Cr16 high-entropy alloy developed via spark plasma sintering

Abstract

This research explores the corrosion resistance and anti-wear property of near-equimolar Ti20-Al16-V16-Fe16-Ni16-Cr16 High Entropy Alloys (HEAs), fabricated at varying temperatures through spark plasma sintering (SPS) procedure. The alloy powders were processed and sintered at temperatures ranging from 700°C to 1100°C, and the resulting HEAs underwent comprehensive characterization. The tribological behavior of these sintered HEAs and their corrosion response in sulfuric and chloride environments were systematically investigated. Microstructural analyses were conducted before and after exposure to corrosive media using a scanning electron microscope. Notably, the assessments in both acid and salt media revealed that HEAs sintered at 1100°C exhibited the highest corrosion resistance, whereas those sintered at 700°C displayed the least resistance. The wear rate is improved by 130% over Inconel 718 alloy in the created HEA. This research signifies the potential of Ti20-Al16-V16-Fe16-Ni16-Cr16 HEA prepared through tailored SPS procedures, as corrosion and wear resistance alternatives surpass the performance of Ti6AlV4 and traditional alloys.

Keywords

high entropy alloys spark plasma sintering potentiodynamic polarization wear resistance and coefficient of friction

Introduction

The continual progress in engineering materials has ushered in a new era of metal alloys, supplementing the existing repertoire of applications. A noteworthy addition to this domain is the emergence of high entropy alloys (HEAs), distinguished by their equimolar compositions comprising a minimum of five elemental metals. These alloys showcase remarkable properties, encompassing elevated corrosion resistance,^1–4 commendable wear resistance,^5,6 favorable electrical and magnetic attributes,⁷ as well as high fracture toughness and fatigue resistance,^7,8 among numerous others. The distinctive characteristics of HEAs are rooted in the cooperative synergy among multiple elements, adopting structures such as body-centered cubic (BCC), face-centered cubic (FCC), hexagonal close-packing (HCP), or their intricate combinations within HEAs rather than generating various intermetallic phases.^1,9,10 Unlike traditional alloys, which grapple with the complexities of developing intermetallic compounds that might compromise mechanical properties, HEAs circumvent these limitations. Their high configurational entropy and sluggish diffusion tendencies facilitate the formation of straightforward crystal-structured solid-solution phases,^11–13 steering clear of the intricate challenges associated with traditional alloy development. This marks a significant departure from the conventional constraints, positioning HEAs as innovative materials with a promising future in diverse engineering applications.

Researchers have used a variety of fabrication techniques to create HEA, including; severe decomposition,¹⁴ ultrasonication,¹⁵ carbothermal shock synthesis,¹⁶ and fast-moving bed pyrolysis.¹⁷ However, it was found that these techniques were either not economically feasible or had low productivity, which called for additional research in the field to increase production efficiency and industrial scalability.¹⁸ The benefits of the process led to the selection of spark plasma sintering (SPS) for the development of HEA. As a result of the use of pulsed DC during SPS, which lowers the activation energy of defect diffusion and migration, impurities and voids are evaporated, leading to reduced porosity in HEAs formed using SPS than those made using the hot pressed (HP) approach, according to Fu et al.¹⁹ As a result, SPS is one of the best methods for creating HEA with little grain development, and interlocked grain boundaries.^20–22 By using mechanical alloying (MA) and SPS to create Co0.5FeNiCrTi0.5 HEA, Fu et al.²³ significantly improved the microstructure and microhardness of the material. NiCoCrCuFe HEA produced by SPS has superior mechanical and microstructural qualities compared to that produced via stir casting, according to Yeh et al.²⁴. When Cr is added to HEA systems, Qiu et al.²⁵ found that the systems’ resistance to corrosion and pitting is increased.

Due to its low density, high strength, high modulus, and low coefficient of thermal expansion (CTE), Ti6Al4 V has gained significant attention in recent years for use in high temperature/strength applications in the aerospace, medical, turbine, chemical storage, and marine industries.^18,26 However, its weak shear strength and corrosive acid sensitivity in HCl or H₂SO₄ limit its industrial applications. The creation of Ti20-Al16-V16-Fe16-Ni16-Cr16 HEA with SPS was envisioned to overcome these Ti6Al4V's shortcomings. The inclusion of ferromagnetic properties of Ni and Fe increase modulus, wear resistance, resistance to oxidation, and resistance to corrosion in an acidic medium. Ni also has an FCC crystal structure, and during HEAs formation, FCC crystal structures cause a higher lattice distortion effect, which reduces strain hardening and prevents crystal dislocation. Contrarily, Fe is an α-phase BCC crystal at ambient temperature, but when heated to 910°C, it changes to an FCC γ-phase, which will help to increase strength and solid solution evolution. To facilitate the evolution of the BCC phase, which increases the ductility and corrosion resistance of the HEAs, Cr was added to the alloys.^27–29 The inherent properties of the Ti-Al-V alloy are notably enhanced by the inclusion of chromium (Cr), nickel (Ni), and iron (Fe), imparting superior characteristics in terms of oxidation resistance, corrosion resistance, and creep resistance.

Also, tribological characterization of HEA has drawn attention from various researchers recently. The BCC phase of HEA is the most desired phase for enhancement of wear characteristics of a HEA. This is because, despite the fact that flexibility is hindered, the BCC phase promotes resistance to brittleness, which improves anti-wear feature. Hence, this study developed HEAs with capabilities surpassing those of the conventional TiAl4 V alloy, particularly when subjected to corrosive and wear environments. The SPS procedure was chosen as the fabrication method, offering a versatile and controlled process for synthesizing advanced materials. The key focus lies in exploring the impact of different sintering temperatures during the SPS process on the resulting HEAs, with the anticipation that varying temperature conditions will yield distinct microstructures and, consequently diverse properties.

Experimental method

Materials

In the experimental setup, a diverse array of powders was employed, each distinguished by specific attributes. The titanium (Ti) powder, boasting 99.9% purity, particle sizes spanning 45–55 µm, and a density of 4.54 gcm⁻³, played a pivotal role. Similarly, the aluminum (Al) powder, with a purity of 99.9%, a particle size of 45 µm, and a density of 2.7 gcm⁻³, contributed to the composition. The vanadium (V) powder, featuring 99.9% purity, particle sizes ranging from 45–50 µm, and a density of 6.1 gcm⁻³, was also incorporated into the experimental matrix. Furthermore, the iron (Fe) powder, boasting 99.9% purity, a particle size of 45 µm, and a density of 7.87 gcm⁻³, played a crucial role in the experiment. Nickel (Ni) powder, known for its 99.9% purity, particle sizes ranging from 45–65 µm, and density of 8.9 gcm⁻³, was an integral component. Finally, the chromium (Cr) powder, characterized by 99.9% purity, particle sizes between 40–65 µm, and a density of 7.19 gcm⁻³, added another dimension to the experimental mix. All these powder materials were sourced from TLS Technik in Germany.

Method

The fabrication process involved the precise blending of these powders in a tubular mixer, with each component weighed accurately to achieve a near equiatomic HEA composition of Ti20-Al16-V16-Fe16-Ni16-Cr16. Ensuring a consistent dispersion of particles, the mixing persisted for 10 h at a speed of 69 revolutions per minute. To enhance homogeneity, steel balls with an 8-mm diameter were introduced to the powder mix at a ball-to-powder ratio of 4:5.³⁰ The precise amount of material required to form a pellet with a 40-mm diameter and a 5-mm thickness was carefully measured and deposited into a graphite die. The prepared mixture of powders was then introduced into the sintering chamber of the SPS machine (KCE-FCT-HHPD 25, Germany). Subsequently, the machine was initialized with specific pressures: vacuum pressure set at 70.5 Pascal, relative pressure at 69 $\times$ 10⁴ Pascal, and absolute pressure maintained at 100 Pa. Once the machine was successfully started, the initial sintering settings for the first sample were configured as follows: a temperature of 700°C, pressure set to 40 MPa, dwell time of 10 min, and a heating rate of 100°C. Subsequently, the sintering temperature was systematically increased to 800, 900, 1000, and 1100°C while keeping the other fabrication parameters constant. After the completion of each sintering cycle, the pellet underwent a thorough sandblasting process to eliminate contaminants from the graphite paper shield.³¹ The resultant set comprised five distinct samples denoted as 700, 800, 900, 1000, and 1100°C, respectively. Following this, the specimens were prepared for metallography, with subsequent division into different sizes for microstructural analysis, corrosion and wear testing. The samples were subjected to etching in de-ionized water using Kroll's reagent and potassium hydroxide. Also, polishing the cut samples to a mirror surface using 120-grit emery cloth lubricated with fumed 0.1 µm silica, ensuring optimal conditions for subsequent analyses was conducted.

Characterization and measurement

With the aid of an X-ray diffractometer (XRD—Bruker D8, Madison, WI, USA), a scanning electron microscope (SEM) coupled to an electron dispersion spectroscope (SEM-EDS, TESCAN), the crystal structure and microstructure of the materials were examined. Using the rule of mixture, the theoretical density (bulk density) of the combined powders was determined using Equation 1.

ρ_{b u l k} = \frac{1}{\frac{A_{w t . %}}{ρ_{A}} + \frac{B_{w t . %}}{ρ_{B}} + \dots + \frac{Z_{w t . %}}{ρ_{Z}}},

(1)

Where

ρ_{A}

is the density of element A,

ρ_{B}

is the density of element B, and

ρ_{Z}

is the density of the last element Z,

ρ_{b u l k}

is theoretical or bulk density (g/cm³), A, B… and Z are the constituent powders that made up the alloy.

The Archimedes approach was used to compute the experimental density.³² Equation 2 was used to compute the relative densification (R.D) of the sintered samples.

R . D = \frac{ρ_{e x p e r i m e n t}}{ρ_{t h e r o t i c a l}} \times 100 %

(2)

where

ρ_{e x p e r i m e n t}

is density obtained from an experiment,

ρ_{t h e r o t i c a l}

is bulk or theoretical density. Equation 3 was used to calculate the porosity of the sintered samples.

P o r o s i t y = 100 % - R e l a t i v e d e n s i t y

(3)

Polarization tests were conducted using NOVA software 1.10 in conjunction with an Autolab Potentiostat (PGSTAT302N). The potentiodynamic polarization experiment employed start and stop potentials of −1.5 V and +1.5 V, respectively, with a precise scan rate of 0.01 V/s. In this setup, the test sample functioned as the working electrode, platinum as the counter electrode, and an Ag/AgCl electrode immersed in a KCl solution served as the reference electrode. Two distinct electrolytes were employed for the polarization tests: a 3.5% NaCl solution and a 0.1 M H₂SO₄ acid solution. Notably, the Tafel extrapolation method was employed to scrutinize the corrosion profile, facilitating the calculation of pivotal parameters such as polarization resistance, corrosion rate, and other noteworthy properties. Tribometer (Anton Paar) version 8.1.8 was used to conduct the tribology test on the samples, which were subjected to loadings of 10 N and 20 N respectively. The counterface was a stainless-steel ball with a diameter of 6 mm. The tribometer was programmed at a duration of 900 s, a frequency of 15 Hz, and a linear speed of 0.06 cm/s. Both the dry and the wet (deionized water) phases of the experiment were carried out at room temperature. The wear rate was directly measured using a profilometer (Surtronic S128) that is connected to the tribometer. This procedure ascertained the rate of wear by measuring the profiles of the wear tracks resulting from plastic deformation. This study presents the mean values of three experiments conducted to measure the wear response. The wear tracks on the samples were viewed and captured using an optical microscope attached to a nanoindenter (Anton Paar, NHT). The optical microscope has an attached personal computer for parameters’ input, image display, and capturing.

Results and discussion

Illustrated in Figure 1 are the SEM micrographs depicting the as-received powders, which reveal prevalent spherical particle morphology. A notable aspect to be highlighted is the electrochemical investigation is the further demonstration of how these spherical particle components contributed to the enhancement of the corrosion resistance of the spark plasma sintered HEAs. The utilization of spherical grains in constituent powders is associated with both a high relative density and a comparatively slower grain development when compared to particles of alternative shapes.³³ This characteristic adds a layer of significance to the role of particles’ morphology in the fabrication process and materials’ properties.

Figure 1.

SEM micrographs of powders (a) Ti, (b) Al, (c) V, (d) Fe, (e) Ni and (f) Cr.

Figure 2 presents the X-ray diffraction graph depicting the received powders. Among all the constituent elements, titanium (Ti) stands out with the highest number of diffraction peaks. The particle sizes in this context are in microns rather than nanometers. This distinction is crucial as micron-sized particles typically exhibit thin and well-defined peaks, in contrast to the strained broad peaks often associated with nano-sized particles.³⁴ The detailed analysis of these X-ray diffraction patterns provides valuable insights into the crystalline structure and composition of the powders, particularly emphasizing the distinctive characteristics attributed to the particle sizes in the micron range.

Figure 2.

XRD of the powders.

Figure 3 showcased the SEM images of the sintered High HEA samples. In Figure 3(a), the presence of porosity in the sample sintered at 700°C indicates insufficient temperature for full alloy densification. This outcome suggests incomplete impurity vaporization and particle micro-diffusion, factors influenced by the pulsed current in Spark Plasma Sintering.²² This deficiency in the sintering process is likely responsible for the observed weak corrosion resistance, a characteristic that will be further explored in subsequent sections. For Figure 3(b), which represents a sample sintered at 800°C showed improvements. The number of pores was decreased, and distinct phases were evolved. However, a few remaining pores and weak grain boundaries persist. As the sintering temperature increases, there is a noticeable enhancement in grain boundary strengthening and compaction. In Figure 3(c), the sample sintered at 900°C exhibits complete crystallization, with additional compaction and densification contributing to the nearly closed pores. Although some dark spots indicative of flaws is still visible, the overall characteristics of this HEA sample have significantly improved. Figure 3(d) presents a sample sintered at 1000°C, revealing a microstructure with two separate crystallized phases. At this temperature, impurity vaporization, grain boundary compaction, and micro-diffusion of particles contribute to maximal densification under optimal SPS conditions. The dark phase demonstrates a BCC structure rich in α and β-Ti, Al, and Cr. Meanwhile, the less intense phase comprises of V, Cr, Fe, and Al-rich FCC structures.

Figure 3.

SEM micrographs of HEA sintered at (a) 700°C (b) 800°C (c) 900°C (d) 1000°C (e) 1100°C and (f) EDS of Ti-rich phase.

At sintering temperatures of 700 and 800°C and 900, 1000, and 1100°C, respectively, the alloy's α-stabilizer (Al) and β-stabilizer (V) encouraged the production of α -Ti-rich BCC¹ phase and β-Ti-rich BCC² phase, respectively. Therefore, as the sintering temperature exceeded the temperature at which α-Ti transforms to β-Ti (883°C), V, α β-stabilizer, encouraged the creation of β -Ti-rich BCC². Because of this, the diffraction peak's position changed from 45° in the 700°C and 800°C samples to 41° in the other samples (Figure 4). The lighter FCC phase is rich in V, Cr, Fe, and Al, whereas the darker BCC phase is rich in Ti, Al, and Cr.¹⁸ The corrosion resistance property of the alloy sintered at 1000°C and 1100°C improved due to the evolution of Ti's BCC² phase at higher sintering temperatures.

Figure 4.

XRD patterns of the sintered powders at different sintering temperatures.

Additionally, the introduction of aluminum into the alloy instigated a strengthening process known as lattice distortion, concurrently enhancing the alloy's ductility. The relatively large atomic radius of aluminum facilitated lattice deformation in the crystal, promoting the creation of the BCC phase.³⁵ As detailed in Table 1, the elevation of the sintering temperature corresponded to the improved compaction and relative densification. The XRD analysis in Figure 4 revealed the absence of detrimental intermetallic compounds, attributable to the high entropy effect. This effect facilitated the formation of a solid solution alloy rather than complex intermetallic compounds.³⁶

Table 1.

Relative densification and porosity of sintered Ti20-Al16-V16-Fe16-Ni16-Cr16 HEA.

Sintering temp (°c)	$ρ_{b u l k}$ (g/cm³)	$ρ_{e x p}$ (g/cm³)	Relative densification (%)	Porosity (%)
700	5.26	4.98	94.72	5.28
800	5.26	5.09	96.84	3.16
900	5.26	5.12	97.34	2.66
1000	5.26	5.20	98.86	1.14
1100	5.26	5.20	98.89	1.11

The presence of nickel with its 10 valence electrons surpassing the threshold of 8, played a crucial role in fostering the formation of the FCC phase,³⁷ particularly evident at the diffraction angle of 77° as observed by Luo et al.³⁸ Furthermore, the transition of BCC α-Fe to FCC γ-Fe was observed at a diffraction angle of 82° when the sintering temperature increased from 900 to 1000°C, considering the curie temperature of pure iron (Fe) is 910°C.³⁹ At 1000°C, the high-entropy effect led to the dissolution of constituent elements into a solid solution, depicted in Figure 4 of the XRD, where all elements coexisted in both FCC and BCC phases within a single peak. Although smaller diffraction peaks were present at 1000°C, their crystallinity diminished when the temperature was further elevated to 1100°C. The reduced strength observed at 1100°C can be attributed to the loss of crystal peaks.¹⁸ The predominant diffraction peaks were BCC when the sintering temperature was below 1000°C, but at 1000°C, the α-Fe BCC transformed into the γ-Fe FCC phase. Additionally, consistent with another study,⁴⁰ the transformation from α-Ti BCC¹ to β-Ti BCC² was observed. At a diffraction angle of 42° was the superimposition of BCC and FCC phases, highlighting the intricate phase transformations occurring within the HEA.

Electrochemical analysis

The samples underwent electrochemical tests in aggressive H₂SO₄ and NaCl environments at room temperature, and the potential against log current was plotted in Figure 5. The Tafel parameters were extrapolated from the Potentiodynamic Polarization (PDP) curves in both anodic and cathodic polarization, and the results are summarized in Table 2. It is apparent that all samples exhibit similar curves, suggesting relatively close corrosion properties in the studied media. Insights from PDP studies highlight the materials’ tendency to resist corrosion, showcasing a correlation between a decrease in corrosion current and an increase in polarization resistance and corrosion voltage.^1,41,42 An increase in sintering temperature led to a decrease in corrosion current from 0.021099A and 0.000449A to 0.000157A and 0.0000305A, while polarization resistance increased from 17.322Ω and 585.28Ω to 158.25Ω and 2558.2Ω for H₂SO₄ and NaCl media when comparing samples sintered at 700°C and 1100°C respectively (Figure 5 and Table 2). The PDP and Tafel parameters results reveal that the sample sintered at 1100°C exhibited the highest resistance to corrosion in both sulfuric and chloride media. When comparing the HEA submerged in a sulfuric medium, it shows a lower corrosion current and a higher corrosion potential in chloride medium. This phenomenon is attributed to the persistent passive film on the HEA in chloride, providing more effective shielding than the passive film on the HEA in sulfuric medium. Elevating the sintering temperature resulted in robust intermetallic bonding and adhesion, which effectively shielded the Ti20-Al16-V16-Fe16-Ni16-Cr16 against potential attack by chloride Cl⁻ ions and sulfate $S O_{4}^{2 -}$ in saline and acidic environments. This suggests that dislocation is prevented by raising the sintering temperature from 700 to 1000, and 1100°C, significantly mitigating the corrosion of the Ti20-Al16-V16-Fe16-Ni16-Cr16 alloy. A similar observation was reported by Ogunbiyi et al.⁴³

Figure 5.

Potentiodynamic polarization of sintered HEA in (a) H₂SO₄ and (b) NaCl.

Table 2.

TAFEL parameters derived from potentiodynamic polarization curves by linear fit for acid and salt media.

Corrosion Media	Sintering Temperature (°C)	Ecorr (V)	Icorr (A)	Corrosion rate (mm/yr)	Polarisation Resistance (Ω)
H₂SO₄
	700	−0.53577	0.021099	2.4517	17.322
	800	−0.5148	0.007453	1.087	29.466
	900	−0.50322	0.000935	0.86602	80.31
	1000	−0.45329	0.000239	0.27754	147.67
	1100	−0.47688	0.000157	0.18253	158.25
NaCl
	700	−0.98938	0.000449	0.52177	585.28
	800	−0.66116	9.92 × 10⁻⁵	0.11524	1119.7
	900	−0.56153	7.24 × 10⁻⁵	0.084074	1430.6
	1000	−0.49693	4.88 × 10⁻⁵	0.056663	1580
	1100	−0.49222	3.05 × 10⁻⁵	0.035468	2558.2

In a general context, the corrosion initiation in HEAs typically originates from a physical or chemical heterogeneity, encompassing imperfections, intermetallic regions, mechanically fractured areas, grain boundaries, inclusions, or disruptions.^44,45 Figures 6 and 7 present SEM imagery of the sintered HEAs after exposure to aggressive attacks in sulfuric and chloride environments, respectively. A detailed analysis of the SEM images in Figure 6 reveals an enhancement in particles homogeneity with an increase in sintering temperature. However, Figure 6(a) displays pronounced non-uniformity with visible pores and cracks, creating pathways for liquid H₂SO₄ infiltration in the sample sintered at 700°C. This susceptibility to corrosion is reflected in the exceptionally high corrosion rate observed in comparison to samples sintered under different conditions.

Figure 6.

SEM images of HEA at (a) 700°C (b) 800°C (c) 900°C (d) 1000°C (e) 1100°C and (f) EDS of Ti-rich phase after exposure in H₂SO₄ Medium.

Figure 7.

SEM images of HEA at (a) 700°C (b) 800°C (c) 900°C (d) 1000°C (e) 1100°C and (f) EDS of Ti-rich phase after exposure in NaCl medium.

The decrease in micropores, as indicated in Table 1, aligns with the observable features of the SEM images in both sulfuric and chloride mediums, indicating that higher sintering temperatures contribute to reduced porosity. Samples sintered at 1000°C and 1100°C exhibit superior corrosion resistance, evidenced by their fine and smooth surfaces, fewer pores, and microcracks (Figures 6(d) and 6(e)). This trend suggests that elevated sintering temperatures promote microstructural uniformity in the SPSed Ti20-Al16-V16-Fe16-Ni16-Cr16 alloys, limiting plausible localized corrosions caused by galvanic coupling at 1000°C and 1100°C during exposure to H₂SO₄ and NaCl solutions. The high densification observed in these samples is also correlated with their remarkable corrosion resistance. Comparing the attacks on sintered samples in different mediums, those exposed in the NaCl medium exhibit superior resistance to corrosion compared to those in the H₂SO₄ medium, as evident in their corroded SEM images (Figures 6 and 7). This correlation aligns with their Tafel parameters detailed in Table 2. It is plausible to assume that metallurgical parameters, including alloy composition, microstructure, and elemental partitioning to phases, play a regulatory role on how these alloys respond to various environmental circumstances.⁴⁶

Tribological analysis

The resulting HEA samples’ coefficient of friction (COF) measured in both dry and wet tribology modes is shown in Figure 8, which was measured at a steady state and presented as the average of the overall wear test duration as recorded by the profilometer (Surtronic S128). The COF shows that the COF reduces as the sintering temperature increases, reaching its lowest level at 1000°C before beginning to rise once more. This pattern demonstrates that as the material's metallurgical and thermodynamic stability rises, so does its COF. It was also demonstrated that the COF reduces as the applied load increases. This is well demonstrated by Amonton's law of friction, which indicates that as the applied load increases, the COF of metallic materials falls.⁴⁷ As a result, increasing the input load causes much heat to be created, which raises the material's temperature and sets off reactions that weaken the material's thermodynamic, intermolecular, and metallurgical bonding.⁴⁸ Soft material experiences less COF than hard and brittle material because it acts as a lubricating agent. Additionally, it was noted in Figure 8 that COF in the dry condition was larger than COF in the wet condition across all samples. This is because deionized water reduces the material's COF by acting as a lubricant and a coolant.⁴⁹

Figure 8.

Coefficient of friction of samples tested on dry and wet modes.

With a COF of 0.025 for the sample sintered at 1000°C and a COF of 0.254 for the sample sintered at 700°C, the sample sintered at 1000°C had a 916% higher COF than the sample sintered at 700°C. When the sample was sintered at lower temperatures, the uncrystallized BCC and FCC phases and insufficient metallurgical bonding were the key contributors to the sample's increased COF. This is due to yet to be developed crystalline BCC and FCC phases which enhance ductility, hardness, modulus, and resistance to galling and oxidation⁵⁰ in HEA and afterwards induce good tribological characteristics, as revealed by a sample sintered at 1000°C. The availability of adequate components helped the HEA's improved wear resistance even more. This is because research has revealed that Cr⁵¹ is a crucial element known for enhancing the tribological properties of materials, while V⁵² and Fe⁵³ create critical oxides that enhance the anti-wear of HEA even at high temperatures. Under applied stresses of 10 N and 20 N, Figure 9 displays the tribology rate of samples placed through wear testing in both dry and wet conditions, which shows how the wear resistance increased as the sintering temperature increased. In general, the addition of Al to the alloy enhanced the wear characteristics of the produced HEA, which also increased the crystallization of the BCC phase,⁵⁴ and increasing strength.

Figure 9.

Wear rate of sintered samples of HEA.

Materials with a large mechanical strength have a high anti-wear, according to Chuang et al.⁵⁵ A material's resistance to erosion and indentation also affects how resistant it is to wear, according to Archard.⁵⁶ It can be proven that in both dry and wet conditions, the wear rate increased as the applied stress increases. According to the observations made by other authors, higher loads result in greater impact and shear forces, which in turn causes more materials to delaminate.^57,58 The quantity of material removed per unit load decreased when water acted as a coolant to decrease heat,⁵⁹ hence the tribology rate measured in the dry mode was greater than the tribology rate recorded in the wet condition. At 20 N of loading (dry mode), the sample fabricated at 700°C exhibited the maximum wear rate of 8.2 × 10⁻³ mm³N⁻¹m⁻¹, whereas the sample fabricated at 1000°C exhibited the lowest wear rate of 1.3 ×10⁻⁴ mm³N⁻¹m⁻¹. This value generally shows that it is better than the tribology rates of a number of popular alloys, such as Inconel 718, which has a value of 3–6 ×10⁻⁴ mm³N⁻¹m⁻¹ (an enhancement of 130%), AISI 316L steel, which has a wear rate of approximately 1 ×10⁻³ mm³N⁻¹m⁻¹ (an improvement of 666%), and TiAl and Fe₃Al alloys, which have wear rates of approximately 1 ×10⁻³ mm³N⁻¹m⁻¹ (same as 666%).^60–62 Notably, raising the fabricating temperature from 700°C to 1100°C improved the wear rate by 620% for similar sample.

Figure 10 displays the COF patterns of the samples as they underwent a dry condition wear test. The two primary stages of the profiles are run-in and steady-state stages.⁶³ Friction force begins by rapidly increasing from zero to its steady state. The amount it moves up depends on the material's strength, asperity, and applied load.^64,65 Figure 10(a) displays the views at 10 N loading. The sample fabricated at 1000°C had a COF that increased to 0.132, and then plummeted to 0.01 before remaining at a steady state of 0.036, whereas the sample sintered at 700°C had the maximum COF range, which was about 0.17 to 0.51. This suggests that the earlier sample was more metallurgical stable than the latter. Also, Figure 10(b) shows that the COF range for sample sintered at 1100°C is around 0.06 to 0.12, at 1000°C is around 0.12 to 0.12 and at 700°C is around 0.17 to 0.34. Therefore, it can be acknowledged that as applied force increases, the COF decreases. This is due to the fact that increased force produces increased heat, which softens the material and lowers COF.⁶⁵

Figure 10.

Coefficient of friction profiles of sample tested on dry condition at (a) 10 N (b) 20 N.

Figure 11 displays the COF patterns for samples examined using deionized water in the wet condition. The COF dropped as the applied forced rose in accordance with Amonton's law.⁴⁷ Water encourages the development of Al, V, and Fe oxides that can act as lubricating films more than in a dry medium.^53,66,67 As can be shown, the alloys’ thermodynamic and metallurgical stability may be predicted using the COF profiles. Samples sintered at higher temperatures (1000°C and 1100°C) displayed more stable profiles compared to samples sintered at lower temperatures. Figure 12 depicts the wear scars left on the samples that were put through a dry mode tribology test with a 20 N loading. As the fabrication temperature increased from 700°C to 1100°C, the tribological condition changed from adhesive only to adhesive and delamination to finally abrasive tribology.

Figure 11.

Coefficient of friction profiles of sample tested on wet mode at (a) 10 N and (b) 20 N.

Figure 12.

Wear track profiles of samples subjected to tribological tests (a) 700°C (b) 800°C (c) 900°C (d) 1000°C and (e) 1100°C.

Figure 12(a) depicts the tribology track on a sample fabricated at 700°C. As whitish material was being removed, it was obvious that the type of wear was adhesive wear. Material erosion left behind grooves that run parallel to the sample's direction of motion. The sample showed weak metallurgical bonding, as evidenced by the significant amount of components that had been removed. Figure 12(b) displays the wear track on a sample that was sintered at 800°C. It is clear that materials were removed, however not as much as the sample sintered at 700°C. From a mechanical and tribological standpoint, this sample is preferable to a sample sintered at 700°C. The wear track for the sample sintered at 900°C is displayed in Figure 12(c), which shows adhesive and delamination wear mode. There was a lot of wear debris in this sample because so many materials were removed. This demonstrates the sample's weak bonding and weak intermolecular forces. In Figure 12(d), the sample that was sintered at 1000°C is depicted, which reveals that the material pull off was significantly decreased and the wear mode had switched to an abrasive one. Despite their short depth, the grooves run parallel to the direction of motion. Figure 12(e) displays a sample that was sintered at 1100°C. Abrasive wear mode and white wear debris were the signs of wear.

This study has investigated the tribological and corrosion properties of Ti20-Al16-V16-Fe16-Ni16-Cr HEA sintered at different temperatures. However, it is recommended that further studies should be done to bridge the limitations of this present study. Further study should be on similar or different HEA formulations using the same spark plasma sintering technology but considering other sintering parameters such as input power, sintering duration, holding time, and applied pressure. Another advanced fabrication method with parameters’ optimization such as laser additive manufacturing can also be used on similar or different HEA formulations. With such HEA development, other properties determination should be focused such as hot corrosion, triobocorrosion, mechanical properties, and thermal stability.

Conclusion

The SPS development of Ti20-Al16-V16-Fe16-Ni16-Cr HEA yields the following findings:

Optimal SPS temperatures of 1000°C and 1100°C result in samples characterized by a fully blended solid solution of HEAs and refined particles rich in both FCC and BCC phases.

Among the samples, those sintered at 1100°C demonstrate the highest corrosion resistance, while those sintered at 700°C exhibit the least. The full development of a solid solution at 1100°C coupled with subsequent stabilization, is credited with enhancing corrosion resistance.

The developed HEA displays minimal corrosion rates of 0.18253 mm/yr and 0.035468 mm/yr when subjected to acidic and salty media at room temperature. This signifies a remarkable 92% and 94% reduction compared to HEA at 700°C.

At a loading of 20 N, the developed HEA showed the lowest tribology rate of 1.3 × 10⁻⁴ mm³N⁻¹m⁻¹. This amounted to a 130% and 666% improvement over Inconel 718 and TiAl alloy, respectively. The sample fabricated at 1000°C had the lowest COF (dry mode) of 0.084.

The tribology experiments’ dry mode had a higher wear rate than the wet mode because deionized water acted as a coolant and lubricating agent to lessen frictional coefficient. The tribological rate increased as the load increased, although COF reduced as the applied force increased because larger load produces more heat and remove more material, which reduces material properties and decreases COF.

Footnotes

Acknowledgment

The authors would love to appreciate Tshwane University of Technology for their support towards this study.

Authors’ contribution statements

Uwa O. Uyor: Conceptualization, Methodology and Writing Original Draft

Abimbola P. I. Popoola: Validation, Funding acquisition and Review/Editing

Olawale M. Popoola: Supervision, Project administration, Validation.

Chika O. Ujah: Investigation and Methodology

Data availability statement

The authors confirm that the data supporting the findings of this study are available within the article.

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

ORCID iD

Uwa O. Uyor

Funding

The authors received no financial support for the research, authorship, and/or publication of this article.

References

Aliyu

Srivastava

. Microstructure-corrosion property correlation in electrodeposited AlCrFeCoNiCu high entropy alloys-graphene oxide composite coatings. Thin Solid Films 2019; 686: 137434.

Tsai

M-H

Yeh

J-W

. High-entropy alloys: a critical review. Mater Res Lett 2014; 2: 107–123.

Liu

Liaw

. Microstructures and properties of high-entropy alloy films and coatings: a review. Mater Res Lett 2018; 6: 199–229.

Miracle

Senkov

. A critical review of high entropy alloys and related concepts. Acta Mater 2017; 122: 448–511.

Gan

Yan

, et al. A non-equiatomic FeNiCoCr high-entropy alloy with excellent anti-corrosion performance and strength-ductility synergy. Corros Sci 2021; 183: 109341.

Yang

Shi

Zhang

, et al. A novel AlCrFeMoTi high-entropy alloy coating with a high corrosion-resistance in lead-bismuth eutectic alloy. Corros Sci 2021; 187: 109524.

Pickering

Jones

. High-entropy alloys: a critical assessment of their founding principles and future prospects. Int Mater Rev 2016; 61: 183–202.

Tang

Yuan

Tsai

C-W

, et al. Fatigue behavior of a wrought Al0. 5CoCrCuFeNi two-phase high-entropy alloy. Acta Mater 2015; 99: 247–258.

Chen

Yang

Dahmen

, et al. Microstructures and crackling noise of AlxNbTiMoV high entropy alloys. Entropy 2014; 16: 870–884.

10.

Jien-Wei

. Recent progress in high entropy alloys. Ann Chim Sci Mat 2006; 31: 633–648.

11.

Peng

, et al. Remarkable cavitation erosion–corrosion resistance of CoCrFeNiTiMo high-entropy alloy coatings. Corros Sci 2021; 190: 109663.

12.

Chou

H-P

Chang

Y-S

Chen

S-K

, et al. Microstructure, thermophysical and electrical properties in AlxCoCrFeNi (0 ≤ x≤ 2) high-entropy alloys. Mater Sci Eng B 2009; 163: 184–189.

13.

Zhang

Pan

, et al. Microstructure and properties of 6FeNiCoSiCrAlTi high-entropy alloy coating prepared by laser cladding. Appl Surf Sci 2011; 257: 2259–2263.

14.

Bondesgaard

Broge

NLN

Mamakhel

, et al. General solvothermal synthesis method for complete solubility range bimetallic and high-entropy alloy nanocatalysts. Adv Funct Mater 2019; 29: 1905933.

15.

Liu

Zhang

Okejiri

, et al. Entropy-maximized synthesis of multimetallic nanoparticle catalysts via a ultrasonication-assisted wet chemistry method under ambient conditions. Adv Mater Interfaces 2019; 6: 1900015.

16.

Yao

Huang

Xie

, et al. Carbothermal shock synthesis of high-entropy-alloy nanoparticles. Science 2018; 359: 1489–1494.

17.

Gao

Hao

Huang

, et al. Synthesis of high-entropy alloy nanoparticles on supports by the fast moving bed pyrolysis. Nat Commun 2020; 11: 2016.

18.

Ujah

Popoola

, et al. Mechanical and oxidation characteristics of Ti20-Al16-V16-Fe16-Ni16-Cr16 high-entropy alloy developed via spark plasma sintering for high-temperature/strength applications. J Mater Eng Perform 2023; 32: 18–28.

19.

Chen

Wen

, et al. Effects of Co and sintering method on microstructure and mechanical behavior of a high-entropy Al0. 6NiFeCrCo alloy prepared by powder metallurgy. J Alloys and Compd 2015; 646: 175–182.

20.

Ujah

Popoola

, et al. Optimisation of spark plasma sintering parameters of Al-CNTs-Nb nano-composite using Taguchi design of experiment. Int J Adv Manuf Technol 2019; 100: 1563–1573.

21.

Thomson

Jiang

Yao

, et al. Characterization and mechanical testing of alumina-based nanocomposites reinforced with niobium and/or carbon nanotubes fabricated by spark plasma sintering. Acta Mater 2012; 60: 622–632.

22.

Oliver

Sunday

Christain

EI-EI

, et al. Spark plasma sintering of aluminium composites—A review. Int J Adv Manuf Technol 2021; 112: 1819–1839.

23.

Chen

Xiao

, et al. Fabrication and properties of nanocrystalline Co0. 5FeNiCrTi0. 5 high entropy alloy by MA–SPS technique. Mater Des 2013; 44: 535–539.

24.

Yeh

J-W

Lin

S-J

Chin

T-S

, et al. Formation of simple crystal structures in Cu-Co-Ni-Cr-Al-Fe-Ti-V alloys with multiprincipal metallic elements. Metall Mater Trans A 2004; 35: 2533–2536.

25.

Qiu

Gibson

Fraser

, et al. Corrosion characteristics of high entropy alloys. Mater Sci Technol 2015; 31: 1235–1243.

26.

Agapovichev

Sotov

Kokareva

, et al. Possibilities and limitations of titanium alloy additive manufacturing. In: MATEC Web of Conferences, EDP Sciences, 2018, pp.01064.

27.

Wang

Zhang

Davies

, et al. Evolution of microstructure, mechanical and corrosion properties of AlCoCrFeNi high-entropy alloy prepared by direct laser fabrication. J Alloys Compd 2017; 694: 971–981.

28.

Huang

Zhu

, et al. Influence of Cr content on the microstructure and mechanical properties of CrxFeNiCu high entropy alloys. Prog Nat Sci: Mater Int 2020; 30: 239–245.

29.

Colombini

Rosa

Trombi

, et al. High entropy alloys obtained by field assisted powder metallurgy route: SPS and microwave heating. Mater Chem Phy 2018; 210: 78–86.

30.

Popoola

Aigbodion

, et al. Improving tribological and thermal properties of Al alloy using CNTs and Nb nanopowder via SPS for power transmission conductor. Trans Nonferrous Met Soc China 2020; 30: 333–343.

31.

Shongwe

Diouf

Durowoju

, et al. Effect of sintering temperature on the microstructure and mechanical properties of Fe–30% Ni alloys produced by spark plasma sintering. J Alloys Compd 2015; 649: 824–832.

32.

Ujah

Popoola

, et al. Electrical conductivity, mechanical strength and corrosion characteristics of spark plasma sintered Al-Nb nanocomposite. Int J Adv Manuf Technol 2019; 101: 2275–2282.

33.

Miyake

Hirata

Shimonosono

, et al. The effect of particle shape on sintering behavior and compressive strength of porous alumina. Materials 2018; 11: 1137.

34.

Ankem

Greene

. Recent developments in microstructure/property relationships of beta titanium alloys. Mater Sci Eng A 1999; 263: 127–131.

35.

Yuan

Min

Yanxiang

, et al. Microstructure and mechanical performance of AlxCoCrCuFeNi high-entropy alloys. Rare Metal Mater Eng 2009; 38: 1602–1607.

36.

Yeh

Chen

Lin

, et al. Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes. Advanced Engineering Materials 2004; 6: 299–303.

37.

Chen

Qin

Zheng

, et al. Composition design of high entropy alloys using the valence electron concentration to balance strength and ductility. Acta Mater 2018; 144: 129–137.

38.

Luo

Zhao

, et al. Selective laser melting of dual phase AlCrCuFeNix high entropy alloys: formability, heterogeneous microstructures and deformation mechanisms. Addit Manuf 2020; 31: 100925.

39.

. Molecular dynamics simulations of fcc-to-bcc transformation in pure iron: a review. Mater Sci Technol 2017; 33: 822–835.

40.

Afolabi

Popoola

, et al. Temperature effects on microstructure and mechanical properties of sintered high-entropy equiatomic Ti 20 V 20 Al 20 Fe 20 Cr 20 alloy for aero-gear application. Int J Adv Manuf Technol 2020; 108: 3563–3570.

41.

Fan

Chen

Fan

, et al. Alcocrfeni high-entropy alloy coatings prepared by gas tungsten arc cladding: microstructure, mechanical and corrosion properties. Intermetallics 2021; 138: 107337.

42.

Jin

Wang

Lin

, et al. Corrosion resistance and cytocompatibility of tantalum-surface-functionalized biomedical ZK60 Mg alloy. Corros Sci 2017; 114: 45–56.

43.

Ogunbiyi

Jamiru

Sadiku

, et al. Influence of sintering temperature on the corrosion and wear behaviour of spark plasma–sintered Inconel 738LC alloy. Int J Adv Manuf Technol 2019; 104: 4195–4206.

44.

Mishra

Balasubramaniam

Tiwari

. Corrosion inhibition of 6061-SiC by rare earth chlorides. Anti-Corros Methods Mater 2007; 54: 37–46.

45.

Alaneme

Olubambi

. Corrosion and wear behaviour of rice husk ash—Alumina reinforced Al–Mg–Si alloy matrix hybrid composites. J Mater Res Technol 2013; 2: 188–194.

46.

Potgieter

Olubambi

Cornish

, et al. Influence of nickel additions on the corrosion behaviour of low nitrogen 22% cr series duplex stainless steels. Corros Sci 2008; 50: 2572–2579.

47.

Barber

. Multiscale surfaces and Amontons’ law of friction. Tribol Lett 2013; 49: 539–543.

48.

El Shalakany

Kamel

Khattab

, et al. Improved mechanical and tribological properties of A356 reinforced by MWCNTs. Fullerenes, Nanotubes and Carbon Nanostructures 2018; 26: 185–194.

49.

Zhang

Song

, et al. Preparation and tribological properties of surface-capped copper nanoparticle as a water-based lubricant additive. Tribol Lett 2014; 54: 25–33.

50.

Poletti

Fiore

Gili

, et al. Development of a new high entropy alloy for wear resistance: FeCoCrNiW0. 3 and FeCoCrNiW0. 3 + 5 at.% of C. Mater Des 2017; 115: 247–254.

51.

Tian

, et al. Effect of chromium content on microstructure, hardness, and wear resistance of as-cast Fe-Cr-B alloy. J Mater Eng Perfor 2019; 28: 6428–6437.

52.

Alvi

Akhtar

. High temperature tribology of CuMoTaWV high entropy alloy. Wear 2019; 426: 412–419.

53.

Hsu

C-Y

Sheu

T-S

Yeh

J-W

, et al. Effect of iron content on wear behavior of AlCoCrFexMo0. 5Ni high-entropy alloys. Wear 2010; 268: 653–659.

54.

Cheng

Liu

Tang

, et al. Microstructure and mechanical properties of FeCoCrNiMnAlx high-entropy alloys prepared by mechanical alloying and hot-pressed sintering. J Alloys Compd 2019; 775: 742–751.

55.

Chuang

M-H

Tsai

M-H

Wang

W-R

, et al. Microstructure and wear behavior of AlxCo1. 5CrFeNi1. 5Tiy high-entropy alloys. Acta Mater 2011; 59: 6308–6317.

56.

Archard

. Contact and rubbing of flat surfaces. J Appl Phy 1953; 24: 981–988.

57.

Yilmaz

. Comparison on abrasive wear of SiCrFe, CrFeC and Al2O3 reinforced Al2024 MMCs. Tribol Int 2007; 40: 441–452.

58.

Ujah

Popoola

, et al. Modification of Al alloy nanopowder with Nb nanopowder on its thermal and tribological properties with SPS for power conductors. Mater Res Express 2019; 6: 116592.

59.

Guo

Jin

Shi

, et al. Tribological behavior of boronized Fe40Mn20Cr20Ni20 high-entropy alloys. Metals 2021; 11: 1561.

60.

Cheng

, et al. Vacuum tribological properties of a Ti-46Al-2Cr-2Nb intermetallics. J Tribol 2014; 136: 021604.

61.

Niu

Zhang

Yang

. Tribological behaviour of Fe3Al-Ba0. 25Sr0. 75SO4 self-lubricating composites in vacuum and air. Vacuum 2018; 154: 315–321.

62.

Mishina

Hase

. Effect of the adhesion force on the equation of adhesive wear and the generation process of wear elements in adhesive wear of metals. Wear 2019; 432: 202936.

63.

J-M

Lin

S-J

Yeh

J-W

, et al. Adhesive wear behavior of AlxCoCrCuFeNi high-entropy alloys as a function of aluminum content. Wear 2006; 261: 513–519.

64.

Quintelier

De Baets

Samyn

, et al. On the SEM features of glass–polyester composite system subjected to dry sliding wear. Wear 2006; 261: 703–714.

65.

Ujah

Popoola

, et al. Enhanced tribology, thermal and electrical properties of al-CNT composite processed via spark plasma sintering for transmission conductor. J Mater Sci 2019; 54: 14064–14073.

66.

Chen

M-R

Lin

S-J

Yeh

J-W

, et al. Effect of vanadium addition on the microstructure, hardness, and wear resistance of Al0. 5CoCrCuFeNi high-entropy alloy. Metall Mater Trans A 2006; 37: 1363–1369.

67.

Kasar

Scalaro

Menezes

. Tribological properties of high-entropy alloys under dry conditions for a wide temperature range—a review. Materials 2021; 14: 5814.