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Diese Studie untersucht die Auswirkungen der TiB2-Phase und der Kühlrate auf die Mikrostruktur, den Verschleiß und die Korrosionsbeständigkeit von AlCoFeNiTiB-Legierungen mit hoher Entropie. Die Forschung hebt die Bildung von Basalzell-, TiB2- und L21-Phasen in der Legierung unabhängig von der Abkühlgeschwindigkeit hervor. Die Studie zeigt, dass eine Erhöhung der Kühlrate die Korrosionsbeständigkeit, Härte und Verschleißfestigkeit der Legierung erhöht. Die AlCoFeNiTiB-Legierung in Plattenform, die mit einer höheren Abkühlrate hergestellt wurde, wies eine überlegene Korrosionsbeständigkeit mit einer Polarisationsbeständigkeit von 58,6 kΩ cm2, die niedrigste Korrosionsstromdichte von 0,31 μA / cm2 und die niedrigste Korrosionsrate von 0,0076 mm / Jahr in einer 5% igen NaCl-Lösung auf. Die Härte der Legierung verbesserte sich ebenfalls mit höheren Kühlraten und erreichte 685 HV1 für die Plattenform. Die Verschleißfestigkeit wurde mittels Pin-on-Disc-Tests bewertet, die zeigten, dass das Vorhandensein der TiB2-Phase zu gleichmäßigeren Verschleißspuren beitrug. Die Studie kommt zu dem Schluss, dass höhere Kühlraten zu verbesserten mechanischen und Korrosionseigenschaften führen, was AlCoFeNiTiB zu einem vielversprechenden Werkstoff für industrielle Anwendungen macht.
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Abstract
This paper presents the results of research on the structure and properties of the AlCoFeNiTiB high-entropy alloy (HEA) obtained by two casting methods. Microstructural studies confirmed the formation of the BCC, TiB2, and L21 phases in the alloy, both in the as-cast state and in the form of plates. Mössbauer spectroscopy showed that in the AlCoFeNiTiB alloy, the iron atoms were mainly distributed in the BCC structure. The paramagnetic properties of the alloys at room temperature were also observed. SEM images revealed the presence of a spinodal decomopsition in the tested alloy in both the ingot and plate forms. Increasing the cooling rate from the liquid state had a positive effect on the corrosion resistance of the alloy in environments of 3.5 and 5 pct NaCl solution. The AlCoFeNiTiB plate in the more aggressive environment showed the best corrosion resistance (polarization resistance of 58.6 kΩcm2, corrosion current density of 0.31 μA/cm2 and the lowest weight loss of 0.0076 mm/year). The results of the electrochemical impedance spectroscopy (EIS) study indicated the beneficial effect of higher cooling rate on the protective abilities of the oxide and hydroxide layer which forms on the alloys surface. The hardness of the rapidly cooled plates (685 HV1) was slightly higher than the hardness of the more slowly cooled ingots (648 HV1), which is likely due to the refinement of the alloy’s microstructure and, consequently, the occurrence of fine grain strengthening. The AlCoFeNiTiB HEA ingot and plate, as well as the AlCoFeNiTi HEA, were described by the same average value of the friction coefficient (0.69); however, the surface morphology of the wear tracks showed a beneficial effect of the TiB2 phase and the increased cooling rate.
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1 Introduction
In recent years, there has been a continuously increasing interest in high-entropy alloys (HEAs). Studies on these materials focus on alloys containing various elements, such as Fe, Ni, Cr, Co, Cu, Ti, Ni, V, and Al, with concentrations ranging from 5 to 35 pct. Due to their unique composition, HEAs may have special mechanical and physicochemical properties that are impossible to obtain in conventional alloys, which makes them attractive in many areas.[1‐5] HEAs are considered a class of promising materials for industrial applications due to their ability to be mass produced using existing technologies.[6‐8] Many investigations have been conducted to determine the effects of various alloying elements on the properties of HEAs.
Metalloids, including boron, are introduced into HEA to increase the strength and plasticity of the face-centered-cubic (FCC) crystal structure.[9,10] Boron has a small atomic size, so it can exist as an interstitial atom, and the addition of a small number of interstitial atoms to HEA can contribute to significant solid solution strengthening effect.[11] In study[12] was shown that just 30 ppm (weight) of boron doping in FeMnCrCoNi and Fe40Mn40Cr10Co10 (at. pct) HEAs dramatically improved the mechanical properties of the alloys. The yield strength was increased by over 100 pct and the tensile strength by ~ 40 pct, with better ductility, as the boron contributed to increase in the grain boundary cohesion and a reduction in grain size after recrystallization. Similarly, the introduction of a small amount of boron (0.015 at. pct) into the refractory HEA Al0.1CrNbVMo resulted in an improvement in the yield strength and compressive strength by 15 pct, and in the plasticity by 26 pct in relation to the undoped alloy. The improvement in yield strength by the addition of B to this alloy was ascribed to the combined effects of Orowan, dislocation, and interstitial solid solution strengthening by the boron.[13]
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In addition to improving mechanical properties, boron can be successfully used as an alternative to expensive metals such as Ni, Co, and Cr. Various studies have shown that boron segregation leads to changes in the structure and grain boundary properties, which provide a simultaneous increase in strength and plasticity.[14,15] As shown in studies,[16,17] doping with boron at an appropriate concentration is highly desirable when developing alloys because it significantly improves the cohesion of the FeMnCoCr HEA surface through interfacial segregation. However, above a certain critical level of doping, the use of boron can become detrimental, leading to the weakening of grain boundaries through the formation of brittle boron-containing compounds. For example, in the (AlCoFeMnNi)100−xBx (x = 0, 0.5, 1, and 5 at. pct) HEA produced by mechanical alloying, an optimal addition of 0.5 at. pct boron was found. Above this value, a reduction in the maximum shear strength was noted.[18]
Boron is typically added to HEAs to improve their hardness and wear resistance, due to its ability to form a boride phase with alloying elements, which acts as a reinforcing second phase. The wear resistance of the Al0.5CoCrCuFeNiBx alloy at x = 1.0 was approximately 20 times higher than that of the alloy at x = 0.[19] It was found that the addition of B to the Al0.5CoCrCuFeNi HEA increased its Vickers hardness to 736 HV (from 232 HV for the B-free alloy).[20] Similarly, in the Al15Cr15Fe50Ni20−xBx (x = 0, 2, 4, 5, 6, and 8; x values in molar ratio) HEAs, the fracture strength and hardness values showed a tendency to first increase and then decrease with increasing B content. The best mechanical properties were obtained with the alloy with a boron content of x = 5 by fine-grain strengthening, second phase strengthening, and solid solution strengthening.[21] In turn, in study,[22] it was stated that the addition of boron in an amount of x ≤ 0.2 (by molar ratio) caused an improvement in the properties of the AlxCoCrFeBx alloy, while a larger addition caused a deterioration. Chen et al.[22] reported that this was due to the increasing volume of the FCC phase with higher boron content.
The impact of boron or boride additions on the corrosion of HEA has so far not been thoroughly studied and explained in depth. The tests carried out on the corrosion resistance of HEAs containing boron showed variable impacts of this element on the corrosion properties.[23] The influence of boron on the improvement of the soft magnetic properties of the AlCoCrFeNiBx alloy (x = 0, 0.1, 0.25, 0.5, 0.75, and 1.0)[22,24] and the deterioration of the soft magnetism of the FeCoNi1.5CuY0.2Bx (x = 0, 0.2, 0.4, 0.6, 0.8, and 1.0) alloy[25] were both noted.
Although work has been previously performed on the doping of HEAs with interstitial elements, there is still much ambiguity about the effects of introducing boron into HEAs. Hence, in this study, a new AlCoFeNiTiB HEA was designed, and its structure, hardness, wear, and anticorrosion properties were investigated.
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2 Materials and Methods
This paper reports the results of studies conducted on AlCoFeNiTiB HEAs in the form of ingots and plates. The ingots, which had a diameter of 30 mm, were produced by induction melting of the pure elements (99.9 wt pct) using a NG-40 induction generator (Łukasiewicz Research Network, Gliwice, Poland) in a ceramic Al2O3 crucible under an argon atmosphere. Plates of 1 mm thickness were prepared by the casting of remelted ingots in a copper mold (with a cooling rate ~ 103 K/s)[26,27] using argon gas under low pressure (0.06 MPa).
Structural analysis of the ingots and plates was performed by X-ray diffraction (XRD) using a MiniFlex 600 (Rigaku, Tokyo, Japan) apparatus. The X-ray diffractometer was equipped with a copper tube (Cu Kα, λ = 0.15406 nm) and a D/TEX strip detector. The measurements were performed on bulk samples in the Bragg-Brentano geometry over an angular range of 2Θ = 20–90° at room temperature. The microstructure of the ingots and plates was examined using a Phenom ProX (Phenom World BV, Eindhoven, Netherlands) scanning electron microscope (SEM). The mapping of the distribution of chemical elements was performed using energy-dispersive X-ray spectroscopy (EDX).
The 57Fe Mössbauer transmission spectra were recorded at room temperature using an Integrated Mössbauer Spectroscopy Measurement System (designed by Wacław Musiał and Jacek Marzec) and a linear arrangement of a 57Co:Rh source, a multichannel analyser, an absorber, and a detector. The spectrometer was calibrated at room temperature with a 30 μm thick α-Fe foil. Numerical analysis of the Mössbauer spectra was performed using the MossWinn4.0i program. Spectral parameters, such as the isomer shift (IS), quadrupole splitting (QS), magnetic hyperfine field (B), full line width at half maximum, and relative subspectrum area, were determined.
To characterize the crystallization process of the ingots, differential thermal analysis (DTA) was conducted using a NETZSCH STA 449 F3 Jupiter thermal analyzer (Netzsch, Selb, Germany). The DTA curves were obtained at 20 °C/min for heating and cooling under a neutral argon atmosphere.
The corrosion resistance of the studied samples was examined in static 3.5 and 5 pct NaCl solutions at 25 °C using an Autolab PGSTAT302N (Metrohm AG, Herisau, Switzerland) potentiostat equipped with a three-electrode cell. The instrument was controlled using NOVA 1.11 software. A saturated calomel electrode (SCE) was used as the reference electrode, and a platinum rod was used as the counter electrode. The changes in the open-circuit potential (EOCP) were recorded at first. Subsequently, electrochemical impedance spectroscopy (EIS) measurements were performed at open circuit potentials. The impedance data were collected over the frequency range of 10–2–105 Hz, using perturbation signal with an AC amplitude of 5 mV. Polarization measurements were conducted in both corrosion environments, in the range of − 400 to 400 mV, with a scan rate of 1 mV/s. The corrosion parameters, such as the corrosion potential (Ecorr) and corrosion current density (jcorr), were determined using the Tafel extrapolation method, using both cathodic and anodic branches of the polarization curves. The polarization resistance (Rp) was calculated according to equation[28]:
where: βa and βc are the anodic and cathodic Tafel slopes, respectively calculated from the Tafel extrapolation. The corrosion rate (vcorr) were subsequently calculated, using the equation[29]:
in which d is the alloy density [g/cm3], K is the constant equal to 3.27 × 103 and EW is the equivalent weight—corresponding to the mass of metal species reacting with charge of 1 F, and can be calculated employing the equation[29]:
where: ni is the valence of “i” element in the alloy, while fi represents its mass fraction in the alloy and Wi its atomic weight.
Hardness measurements were performed using a Future Tech FM-700 Vickers (Future Tech, Tokyo, Japan) hardness testing instrument with a load of 1000 g for 15 seconds. Tribological tests of the ingots and plates were performed with the pin-on-disc method using a tribometer (CSM Instruments, Peseux, Switzerland). Additionally, to compare the effect of boron and the occurrence of the TiB2 phase on wear resistance, AlCoFeNiTi was also tested in the form of an ingot described in a previous study.[30] The radius of the wear track was 1.5 mm, and the counter sample was a ball made of Al2O3 (d = 6 mm). The linear speed was 0.01 m/s, which corresponds to a sliding speed of 6.67 rad/s. A load of 10 N was applied. The tests were carried out in an air atmosphere at room temperature (25 °C). Relative humidity was not determined or controlled during the measurements. Observations of the surface morphology of the wear tracks after the tribological tests were carried out with a Phenom ProX SEM (Phenom World BV, Eindhoven, The Netherlands).
3 Results and Discussion
3.1 XRD Analysis
The structure of the AlCoFeNiTiB alloy was evaluated by XRD (Figure 1). XRD patterns of the samples, both in the as-cast state and in the form of plates, showed peaks corresponding to the crystalline BCC, TiB2, and L21 phases. This indicates that the solidification conditions did not affect the phase constitution of the investigated alloy. Despite the absence of new phases in the rapidly cooled samples, differences in peak intensity and width can be observed, which may be related to the different degree of refinement of the structures, the size of the crystallites, or the internal stresses.
Fig. 1
XRD patterns of AlCoFeNiTiB HEAs in the form of ingot and plate
As reported in the literature,[22,24,25,30‐33] the addition of boron causes the formation of boride phases and structural transformations depending on its concentration and alloying system. Also, in our case, the formation of new phases was found compared to the boron-free AlCoFeNiTi alloy, in which the (Fe, Co, Ni)2TiAl L21 Heusler phases and the BCC phases were identified.[30] It can, therefore, be concluded that the addition of B changed the phase composition of the AlCoFeNiTiB HEA, while the preparation method, whether as-cast or in the plate form, did not affect it. A change in the phase composition after the boron addition was also observed in AlCoCrFeNiBx and AlFeCoNiBx alloys: when the B content increased, the intensity of diffraction peaks of the BCC phase decreased, while diffraction peaks of the FCC phase appeared.[22,33]
3.2 Mössbauer Measurements
The Mössbauer spectra of the AlCoFeNiTiB HEA in an as-cast state (ingot) and in the plate form are presented in Figure 2, along with the fitted components and their phase assignments and contributions. These spectra were fitted with one quadrupole doublet and several sextets. The doublet constituted the central part of the spectrum. Its hyperfine parameters were IS = 0.10(2) mm/s and QS = 0.17(2) mm/s for the ingot alloy and IS = 0.11(1) mm/s and QS = 0.20(1) mm/s for the plate-type alloy, which are characteristic of a paramagnetic phase rich in Al.[34] Such a signal was interpreted as the result of mixing Al and Fe nanoparticles, which react to form the BCC structure.[34‐36] The paramagnetic properties at room temperature for the bcc-Fe(Al) alloys were also observed for these Ti-modified alloys.[37] In the Mössbauer spectrum, the ingot alloy was also fitted using sextets with hyperfine magnetic fields of 36.4(1), 35.0(1), 33.3(1), 31.6(3), 29.5(2), 26.1(3), 21.4(3), 19.9(4), and 9.2(2) T. The obtained values of isomer shift for these sextets were 0.04(1), 0.04(1), 0.02(1), 0.04(1), 0.06(1), 0.08(2), 0.09(2), 0.16(3), and 0.09(2) mm/s, respectively. The first three, which had a hyperfine magnetic field higher than that of pure iron in the BCC phase (33 T), are characteristic of this structure when Fe is substituted by Co or Ni atoms.[38,39] The hyperfine parameters of the five other sextets characterized the five Fe sites in the DO3 structure of the Fe3Al alloy.[40,41] Components with hyperfine magnetic field values of 30 and 21 T are always present for these Fe-rich alloys and represent two inequivalent Fe sites with either no Al or four Al nearest neighbors.[42] The other three magnetic subspectra corresponded to five, three, and two Al configurations at nearest neighbor positions in this structure. The last sextet, with the smallest hyperfine magnetic field, was probably associated with the B-rich Fe-B phase.[40] However, considering the XRD results, this component may be related to the TiB2 phase, in which Fe atoms substitute some Ti atoms. In the Mössbauer spectrum for the plate alloy, in addition to the quadrupole doublet, there were seven sextets with hyperfine magnetic fields equal to 34.9(2), 32.2(1), 29.5(5), 26.2(3), 22.8(3), 18.0(3), and 12.1(2) T, and the related isomer shift values were 0.07(3), 0.03(2), 0.07(2), 0.08(2), 0.11(3), 0.16(3), and 0.15(3) mm/s, respectively. The contribution of components related to the Fe3Al and Ti(Fe)B2 phases practically did not change. However, the concentration of the magnetic component associated with the bcc-Fe(Co,Ni) phase decreased. At the same time, the share of the doublet associated with bcc-Fe(Al,Ti) increased. This trend possibly indicated the redistribution of atoms inside this structure during the preparation of the alloy in the plate form. Changing the local environment of Fe atoms caused a change in the magnetic properties of the phase.
Fig. 2
Mössbauer spectra of AlCoFeNiTiB HEAs in an as-cast state (a) and in the plate form (b). The experimental points, fitting curves and spectral components corresponding to particular Fe phases in different structural positions (colored lines) are presented together with their relative contributions (Color figure online)
Figure 3 shows SEM images of AlCoFeNiTiB HEA in the form of an ingot and plate at the same magnification. Dark precipitates are titanium borides, substantially larger in case of alloys in the form of ingots [Figure 3(a)]. As a result of rapid cooling, the refinement of the alloy structure occurred and large precipitates were defragmented into smaller ones, which can be clearly seen in Figure 3(b). Figure 4 presents SEM images and EDX maps with different concentrations of chemical elements of AlCoFeNiTiB in the ingot form. It can be observed that the structure of the ingot consisted of a matrix, understood as a continuous area, and globular precipitates. Gan et al.[43] described the SEM structure of AlxCoCrFeNi2.1 alloys as consisting of dendritic and interdendritic areas. A similar structure for the AlCoCrFeNi2.1 alloy was also described by Charkhchian et al.[44] For the analyzed AlCoFeNiTiB alloy in the ingot form in the present study, it could also be assumed that the matrix was the interdendritic phase and the globular precipitates were the dendritic phase. The matrix of the studied alloy was depleted in aluminum, which is visible in the EDX map. Moreover, the matrix was a mixture of phases: the lighter area was Fe-rich, while the darker one was Ti-rich.
Fig. 3
SEM images of AlCoFeNiTiB HEA in the form of an ingot (a) and plate (b)
In the study of Gan et al.[43] the eutectic occurring in the interdendritic area in a structure of the Al0.8CoCrFeNi2.1 alloy was indicated. In the matrix area, the nickel content was not uniform. Globular precipitates with spherical shapes were characterized by a gradient effect, which could be observed by the occurrence of spinodal decomposition at the boundaries. Spinodal decomposition is a phase-separation phenomenon that occurs in a homogeneous, supersaturated phase. The resulting phases as a result of spinodal decomposition differ in chemical composition.[45] Inhomogeneous chemical distribution in the form of a “gradient effect” was also visible in the EDX elemental distribution maps. Inside the precipitates, there was a high content of aluminum and a small amount of iron, while the opposite relationship occurred at the boundaries. Spinodal decomposition in the dendritic area of the AlCoCrCuFeNi HEA was previously described by Singh et al.[46] The SEM image with EDX maps for the AlCoFeNiTiB HEA in the form of plates is presented in Figure 5. The structure of the rapidly solidified plates was characterized by the presence of Al- and Ni-rich dendrites. The interdendritic area was rich in iron, while cobalt was homogeneously dissolved in the structure of the plate. As in the case of the ingots, the presence of TiB in the AlCoFeNiTiB plates was observed; however, titanium was also dissolved in the dendritic and interdendritic areas. Moreover, the AlCoFeNiTiB HEA in the plate form was characterized by the presence of spinodal decomposition in the interdendritic area, visible in Figure 6. Based on the results of the previous study[30] on the AlCoFeNiTi HEA, it can be concluded that in the AlCoFeNiTiB HEA, for both cooling rates, the interdendritic region matrix consisted of the more Fe-rich BCC phase, while the dendritic phase consisted of an Al-rich L21 phase.
Heating and cooling curves were recorded for the studied AlCoFeNiTiB HEA in the ingot form using differential thermal analysis (DTA). Figure 7 presents the DTA curves after heating and cooling at 20 °C/min. During heating, three distinct endothermic events were recorded at temperatures of 1114, 1172, and 1382 °C. It can be assumed that the thermal event at 1114 °C corresponded to the beginning of melting of the Ni3Ti phase based on the Ni-Ti phase equilibrium system.[45] The effect at 1382 °C was probably due to complete transition to the liquid state of the tested alloy. The liquidus temperature of the Ni3Ti phase is approximately 1380 °C.[47] However, for the AlFe3 phase identified by XRD, a transformation occurred up to a temperature of 552 °C (AlFe3 → AlFe), and then over a temperature range from 622 to 1022 °C, the AlFe → α-Fe transformation took place in accordance with the Al-Fe phase equilibrium system.[48] The solidus temperature of the α-Fe phase, depending on the aluminum content, can range from 1310 to 1536 °C.[48] Due to the high melting point of the TiB2 phase (above 3000 °C), no effect was recorded for the TiB2 phase, although, based on EDX maps, it is known that the contribution of this phase was very small.[49] It could, therefore, be assumed that the melting of this phase occurred at a lower temperature. During cooling, one exothermic event was recorded at 1382 °C.
Fig. 7
DTA curves of the AlCoFeNiTiB HEA ingot after heating and cooling at 20 °C/min
The corrosion resistance of the prepared alloy was estimated based on the electrochemical tests carried out in 3.5 and 5 pct NaCl aqueous solutions at 25 °C using the potentiodynamic method. The changes in the open circuit potential over time (a) and polarization curves (b) for samples in the as-cast state and in the plate form are shown in Figures 8 and 9. The quantitative parameters of the open circuit potential, corrosion potential, polarization resistance, corrosion current density and corrosion rate were summarized in the Tables I and II.
Fig. 8
Changes in the open-circuit potential with time (a) and polarization curves (b) for ingot and plate in 3.5 pct NaCl solution at 25 °C
Electrochemical Parameters of the AlCoFeNiTiB HEA in 3.5 Pct NaCl Solution
Sample
EOCP [V] (± 0.01)
Ecorr [V] (±0.01)
Rp [kΩcm2] (±0.1)
jcorr [μA/cm2] (±0.1)
vcorr [mm/year] (±0.001)
Ingot
− 0.409
− 0.415
32.1
0.55
0.0138
Plate
− 0.182
− 0.221
38.5
0.34
0.0085
Table II
Electrochemical Parameters of the AlCoFeNiTiB HEA in 5 Pct NaCl Solution
Sample
EOCP [V] (±0.01)
Ecorr [V] (±0.01)
Rp [kΩcm2] (±0.1)
jcorr [μA/cm2] (±0.1)
vcorr [mm/year] (±0.001)
Ingot
− 0.508
− 0.545
7.8
1.24
0.0308
Plate
− 0.229
− 0.321
58.6
0.31
0.0076
The rapidly solidified AlCoFeNiTiB plate characterize with substantially more noble open-circuit and corrosion potentials compared to the alloy in the form of ingot, in both corrosion environments. The most favorable values of the EOCP and Ecorr of − 0.182 and − 0.221 V, respectively, were obtained in the milder 3.5 pct NaCl solution. Concurrently, use of the higher cooling rate led to the decrease in the corrosion current density and increase in the polarization resistance values, evidencing the beneficial effect of increasing the cooling rate during solidification on the corrosion resistance. The highest value of polarization resistance (58.6 kΩcm2), the lowest corrosion current density (0.31 μA/cm2) and the lowest corrosion rate (0.0076 mm/year) were obtained for the AlCoFeNiTiB plate, measured in a more aggressive 5 pct NaCl environment. In turn, the least favorable corrosion parameters—the lowest value of polarization resistance (7.8 kΩcm2), the highest corrosion current density (1.24 μA/cm2) and corrosion rate (0.0308 mm/year) characterize the as-cast alloy in the same corrosion environment. The positive influence of the increase in cooling rate on the corrosion resistance of the alloys can be related to the microstructure fragmentation and resulting more homogenous elemental distribution. A similar relationship of corrosion behavior to the preparation method was shown for the AlCoFeNiTi alloy without boron. Also, in this case, the alloy in the form of a plate had better values for all the analyzed corrosion parameters.[30] Furthermore, comparison with the alloys described in this work shows also the positive influence of boron addition. In particular, for the as-cast alloys, the introduction of boron contributed to a significant decrease in corrosion current density—by two orders of magnitude, as well as increase in the polarization resistance by one order of magnitude. Concurrently, alloys in this work characterize also with more noble corrosion potential values. Differently, in the study,[50] a negative influence of boron addition on the corrosion resistance of the Al0.3CoFeMnNiBx alloy was observed. As the boron content increased, the current density increased, and the corrosion potential decreased, which was attributed to the presence of interdendritic borides in the boron-containing alloys. Areas of concentrated boride are susceptible to corrosion attack, and galvanic cells are likely to form at the boundaries of the boride precipitates, which accelerate corrosion.[50] Similarly, in paper,[20] it was found that Al0.5CoCrCuFeNiBx alloys (x = 0, 0.2, 0.6, or 1.0) in 1 M H2SO4 were not susceptible to local corrosion; however, as the B content increased, the icorr values of the alloys increased, and the repassivation potential (Erp) decreased, indicating that the corrosion resistance of the Al0.5CoCrCuFeNiBx alloy decreased while the boron concentration increased. In turn, in the FeCrNiCoBx alloy in the form of a coating, the addition of B in a range of 0.5 ≤ x ≤ 1.0 resulted in a decrease in the alloy’s corrosion rate, while increasing the element’s content (x = 1.25) resulted in an increase in the corrosion rate. The variable corrosion resistance is explained by the phase transformation of (Cr,Fe)2B borides with a rhombic structure to a (Fe,Cr)2B, tetragonal structure with increasing boron content.[51]
To further investigate the corrosion behavior of the AlCoFeNiTiB alloys, the electrochemical impedance spectroscopy (EIS) measurements were performed. Figures 10 and 11 show the Nyquist and Bode plots, for the tests conducted in the 3.5 and 5 pct NaCl solutions, respectively. All Nyquist spectra characterize with the presence of unfinished semicircles [Figures 10(a) and 11(a)], with a noticeably larger diameter in the case of alloys in form of plates, indicative for a higher surface film impedance.[52] The Bode phase angle plots show that in both environments rapidly solidified alloys characterize with higher maximum phase angles, maintained in the broader frequency range, indicating higher chemical stability of the oxide and hydroxide layer formed on their surface.[53] In turn, for as-cast alloys, the presence of dual broadening on the Bode phase angle plots allows determining the existence of two time-constants in the system.[54,55] Consequently, the electric equivalent circuit model (EEC) involving the presence of two electroactive layers was used to fit the recorded spectra [Figure 12(a)].[54] The EEC used is made up of solution resistance (Rs), charge transfer resistance of the electrical double layer (Rdl), charge transfer resistance of the surface film (Rf), constant phase element representing the double-layer (CPEdl) and the surface film capacitance (CPEf).[55] To analyze the spectra obtained for alloy in the form of plates, the electric equivalent circuit model shown at Figure 12(b) was utilized. In this case, the EEC involves solution resistance (Rs), surface film resistance (Rf), and constant phase element representing the capacitance of the surface film (CPEf).[56]
Fig. 10
Experimental EIS spectra: Nyquist plots (a), Bode modulus plots (b) and Bode phase angle plots (c) for AlCoFeNiTiB alloys in form of ingot and plate in 3.5 pct NaCl solution, at 25 °C
Experimental EIS spectra: Nyquist plots (a), Bode modulus plots (b) and Bode phase angle plots (c) for AlCoFeNiTiB alloys in form of ingot and plate in 5 pct NaCl solution, at 25 °C
The constant phase element was used instead of the capacitance to account for the frequency dispersion, which can be related to surface inhomogeneity.[57] The following equation can be used to calculate the impedance of the constant phase element impedance[58]:
where: Z0 is the capacitance of the capacitance of the parameter related to the electrode [F cm–2 sn–1]; ω represents the angular frequency (rad/s) and n is the constant phase exponent. The CPE can correspond to the various circuit elements, depending on the n value, when the n = 1 represents the pure capacitance, Warburg impedance for n = 0.5 or can be purely resistive, when n = 0.[59]
The electrochemical parameters fitted are summarized in Tables III and IV, for the measurements carried out in the 3.5 and 5 pct NaCl alloys, respectively. For as-cast alloys, the values of polarization resistance, which reflects the corrosion resistance, can be calculated as \({R}_{\text{p}}={R}_{\text{f}}+{R}_{\text{dl}}\).[59] Confirming the tendency revealed as a result of potentiodynamic polarization measurements, rapidly solidified alloys characterize with significantly higher values of polarization resistance in both environments, evidencing the positive effect of a higher cooling rate on the corrosion resistance. However, the constant phase exponent n below 0.8 suggests the formation of a porous surface film.[60] The corrosion behaviour of the plates seems to be less affected by differences in the concentration of chloride ions, considering the similar characteristics of the oxide and hydroxide layer formed. In turn, alloys in form of ingots exhibit noticeably lower polarization resistance in the more aggressive solution, demonstrating the increased susceptibility of the surface film to aggressive ion penetration.[61] The chloride ions diffuse through the defects in the alloys surface film, leading to its dissolution and the occurrence of pitting corrosion. Higher concentrations of chloride ions can disrupt the oxide layer more effectively, facilitating the occurrence of corrosion processes. Furthermore, a more concentrated NaCl solution characterizes with higher conductivity, which accelerates electrochemical reactions.[62,63] In addition, in the case of cast alloys, significantly higher values of surface film capacitance were obtained in both corrosion environments, relating to an increased accumulation of charge on the surface, which can be attributed to the more developed area of the oxide and hydroxide layer formed.[61]
Table III
The Fitted Electrochemical Parameters for Impedance Data of the AlCoFeNiTiB Alloy in 3.5 Pct NaCl Solution, at 25 °C
Sample
RS [Ωcm2]
CPEf [μΩ−1 cm–2sn]
nf
Rf [kΩcm2]
CPEdl [μΩ−1 cm–2sn]
ndl
Rdl [kΩcm2]
AlCoFeNiTiB Ingot
6.39
192.07
0.71
15.43
56.35
0.78
1.03
AlCoFeNiTiB Plate
7.82
17.37
0.81
20.13
Table IV
The Fitted Electrochemical Parameters for Impedance Data of the AlCoFeNiTiB Alloy in 5 Pct NaCl Solution, at 25 °C
Sample
RS [Ωcm2]
CPEf [μΩ−1 cm–2sn]
nf
Rf [kΩcm2]
CPEdl [μΩ−1 cm–2sn]
ndl
Rdl [kΩcm2]
AlCoFeNiTiB Ingot
6.91
131.04
0.75
6.97
41.88
0.83
0.84
AlCoFeNiTiB Plate
1.97
14.86
0.77
21.14
In the work,[64] describing the corrosion resistance of the CoCrFeNiBx alloy (where x = 0, 0.01, 0.05, 0.1, 0.5) in boron 3.5 pct NaCl solution, the alloying with boron was found detrimental to the stability of the passive film, although the alloys still exhibit relatively good corrosion resistance.[64] In turn, Yue et al.[53] investigated the passivation behavior of AlCoCrFeNiTix (where x = 0, 0.25, 0.5, 0.75, 1.0) laser-cladded coatings in chloride ion environment. It was found that the corrosion resistance of the alloys initially increased with the titanium concentration, reaching maximum for the AlCoCrFeNiTi0.5 alloy, following with its subsequent decrease. The negative effect was attributed to the increased microgalvanic activity associated with the precipitation of the Laves and TiC phase. Concurrently, the EIS study revealed that the AlCoCrFeNiTi0.5 alloy characterize with the most stable and protective passive film, exhibiting the polarization resistance of 27.83 kΩcm2.[53]
The varied corrosion resistance of the investigated AlCoFeNiTiB high entropy alloy in NaCl solutions may also be influenced by the presence of the TiB2 phase, the size, and the number of precipitates. Zhang et al.[65] examined FeCoNiCr HEA reinforced with TiB2 nanoparticles produced by laser powder bed fusion. Based on corrosion tests conducted in 3.5 pct NaCl solution, it can be concluded that the compound reinforced with TiB2 particles was characterized by better corrosion resistance compared to FeCoNiCr HEA. Furthermore, Lou et al.[66] produced the HEA coating which exhibited a nanocomposite structure consisting of TiB2 nanocrystallites embedded in an amorphous TiZrNbTaFe matrix. The authors reported that the HEA Ti15.2Zr0.9Nb1Ta1.2Fe1.1B33.8 coating with TiB2 nanocrystalites exhibited a high hardness and critical adhesion load, as well as good wear and corrosion resistance in a 3.5 pct NaCl solution that equimolar amorphous TiZrNbTaFe HEA coating. Therefore, the presence of the TiB2 phase in the AlCoFeNiTiB HEA has an impact on its corrosion resistance in an environment containing chloride ions. On the one hand, TiB2 has a positive effect on corrosion protection by supporting the formation of a stable passive layer. However, the uneven distribution of the TiB2 phase causes the formation of galvanic microcells at the boundary of the interphase, which leads to a local weakening of the corrosion resistance.
3.6 Hardness
The results of the hardness tests of AlCoFeNiTiB HEA in the ingot and plate form are presented in Figure 13. Twenty-five measurements were performed for each, and the plate gave a higher hardness than the ingot (685 HV1 + /− 21 and 648 HV1 + /− 97, respectively). Changes in the measured values contributed to an increase in the standard deviation, especially in the case of the ingot. The observed effect of the cooling rate on the hardness of the investigated alloys can be explained by the presence of a boride phase, which was probably distributed so unevenly that the hardness measurements were noticeably variable. Also, the increase in the hardness of rapidly cooled plates is related to the refinement of the alloy structure and, as a result, the occurrence of fine grain strengthening. Rapid cooling likely inhibited the intergranular segregation of borides, which is commonly observed in slowly solidified alloys and can lead to localized softening or embrittlement.
Similar observations have been reported in other B-containing high entropy alloys, where the increase in hardness was attributed to the formation of hard borides and solid solution strengthening mechanisms.[24,32,67,68] In particular, it has been reported that borides can create effective barriers to dislocation and, in some cases, accumulate in interdendritic regions, leading to local hardening. In the Al0.2Co1.5CrFeNi1.5Ti0.5Bx alloy (x = 0, 0.15, 0.3, 0.45, 0.6, 0.75, or 0.9, in molar ratio), the microhardness increased significantly by boron doping, and the microhardness of the interdendritic regions was much higher than that of the dendritic regions, in line with the finding that the boride structure accumulated mainly in the interdendritic regions.[32]
3.7 Pin-on-Disc Measurements
In order to evaluate the influence of cooling rate, structure, and the presence of the TiB2 phase on the wear resistance of the AlCoFeNiTiB HEA, pin-on-disc tests were performed using an Al2O3 counter sample. The coefficient of friction as a function of time curves for the AlCoFeNiTiB produced with two cooling rates, as well as for the AlCoFeNiTi ingot investigated in the previous study,[30] are shown in Figure 14 for comparison. The AlCoFeNiTiB HEA in the ingot form was characterized by a more stable course of the curve compared to the plate, which showed greater deviation. The curve for AlCoFeNiTi HEA in the ingot form was characterized by higher values of the friction coefficient up to 1200 seconds, after which a significant decrease occurred. The average friction of coefficient value 0.69(± 0.06) was the same for all the studied samples.
Fig. 14
Pin-on-disc curves of the AlCoFeNiTiB ingot and plate examined in this study and the AlCoFeNiTi ingot for comparison
Yadav et al.[69] studied spark plasma sintered AlCrFeMnV HEA with the addition of Bi and TiB2. The soft Bi phase should provide lubrication by forming a thin layer between the contacting surfaces, while the aim of adding the hard TiB2 phase was to obtain higher strength and increase the load-bearing capacity of the alloy system. According to these authors,[69] the addition of TiB2, due to its high hardness, led to a 85 pct reduction in wear rate compared to base alloy. However, the addition of TiB2 increased the friction coefficient from 0.15 (for Bi-HEA) to 0.35 (10 wt pct TiB2-HEA). The increase in the coefficient of friction was attributed to the occurrence of hard TiB2 particles through which deep grooves were formed on the surface, retaining the dispersoids and increasing the roughness.[69]
The SEM images of the worn surface morphology after pin-on-disc tests are compared in Figure 15 in SE and BSE mode and in Figure 16 with marked wear mechanisms. It can be observed that the worn surfaces, especially visible in the BSE mode as dark tracks, contain smeared contamination. The particles formed by abrasion were loosely attached to the surface, prone to cracking and detached from the wear track. In Figure 16, for the AlCoNiFeTi and AlCoNiFeTiB ingots, the brittle crack mechanism was observed in the form of sharp, regular cracks. Furthermore, groove formation was observed for all the studied alloys, while a small amount of wear debris was visible for the AlCoNiFeTi ingot. Table V shows the results of EDX chemical composition analysis from the points marked in Figure 16. It was observed that the point analysis from dark areas has a high oxygen content, which confirms the occurrence of wear products after pin-on-disc tests. In the areas where no wear products were present, a lower oxygen content was identified.
Fig. 15
SEM images of worn surface morphology after pin-on-disc tests for the AlCoFeNiTi ingot investigated in previous paper[30] (a) and AlCoFeNiTiB HEA in the form of an ingot (b) and a plate (c) in SE and BSE mode
SEM-BSE images of the surface morphology and wear mechanisms after pin-on-disc tests for the AlCoFeNiTi ingot investigated in the previous paper[30] (a) and AlCoFeNiTiB HEA in the form of an ingot (b) and a plate (c) with marked points for EDX analysis
Chemical Composition by EDX Analysis of Marked Points (Fig.16) for AlCoFeNiTi Ingot and AlCoFeNiTiB HEA in the form of an Ingot and a Plate After Pin-On-Disc Tests
Sample
Point
O
Al
Co
Fe
Ni
Ti
AlCoFeNiTi Ingot
1
74.07
6.38
4.97
5.27
4.70
4.61
2
74.26
6.47
5.06
4.65
4.58
4.99
3
64.55
6.68
7.95
5.85
7.24
7.73
4
55.93
10.25
9.03
6.72
9.31
8.76
5
22.42
16.10
14.85
18.44
14.96
13.23
AlCoFeNiTiB Ingot
1
80.25
5.74
2.64
2.72
2.78
5.86
2
72.83
7.16
4.95
4.20
4.88
5.98
3
77.16
6.41
3.72
4.19
3.36
5.16
4
48.19
6.56
12.69
14.06
9.69
8.81
5
31.27
21.72
14.92
8.87
15.75
7.46
AlCoFeNiTiB Plate
1
73.40
6.15
5.82
5.35
5.59
3.69
2
76.37
4.34
3.56
3.66
3.53
8.55
3
73.33
5.24
5.86
6.77
4.83
3.98
4
71.79
6.16
6.08
6.16
5.53
4.28
5
67.00
9.19
7.11
5.33
7.05
4.31
Dada et al.[70] investigated the wear mechanisms of laser-deposited AlCoCrCuFeNi alloy. According to these authors, the scanning speed of the laser beam affected the grain refinement in the structure, which led to the increased strength of the alloy. The high aluminum content stabilized the solid solution BCC structure, which was also responsible for the strengthening mechanism of the alloy, resisting the plastic deformation caused by abrasive wear and improving the wear resistance.[68] Similarly to the present study, Feng et al.[71] also studied the AlCrFeNiV HEA using an Al2O3 ball. According to these authors,[71] the occurrence of deformation and grooves is due to the use of a hard counter sample pressed into the soft alloy surface. According to Yadav et al.[69] the uniform distribution of TiB2 particles (together with Bi) contributed to the improved wear resistance. In our study, the examined alloys showed the same average friction coefficient value; however, SEM surface morphology observations showed that the presence of the TiB2 phase contributed to more uniform wear.
4 Conclusions
To obtain the AlCoFeNiTiB HEA, two cooling rates from the liquid state were used—induction melting of the elements in a ceramic crucible (ingots) and pressure casting in a copper mold (plates). The cooling rate did not affect the phase composition of the alloy: both in the as-cast and plate states, the AlCoFeNiTiB alloy consisted of the BCC, TiB2, and L21 phases. However, an inhomogeneous chemical distribution could be observed in the form of a “gradient effect” visible in the EDX elemental distribution maps. Moreover, the AlCoFeNiTiB HEA in the plate form was characterized by the presence of a spinodal distribution in the interdendritic area. Increasing the cooling rate improved the corrosion resistance of the tested alloys, as evidenced by the values of the parameters EOCP, Ecorr, Rp, jcorr and vcorr. The AlCoFeNiTiB alloy in the plate form in 5 pct NaCl solution showed the best corrosion resistance, exhibiting the highest value of polarization resistance (58.6 kΩcm2), the lowest corrosion current density (0.31 μA/cm2) and the lowest weight loss (0.0076 mm/year). The preparation method also affected the hardness and wear resistance of the alloy. A higher hardness (685 HV1) was obtained for the alloy in the plate form, which may be related to the distribution of the TiB2 phase and fine grain strengthening. The AlCoFeNiTiB HEA ingot and plate, as well as the AlCoFeNiTi HEA ingot, were characterized by the same coefficient of friction value (0.69). However, the samples with the addition of boron indicated a more uniform wear track on the SEM images taken after the pin-on-disc measurements. Moreover, for the AlCoFeNiTiB plate, no brittle crack was observed, compared to its observation for the AlCoFeNiTi and AlCoFeNiTiB ingots. Therefore, a higher cooling rate resulted in improved mechanical and corrosion properties.
Acknowledgments
This work was supported by the National Science Centre of Poland under research project no. 2022/47/B/ST8/02465.
Conflict of interest
All the authors of the article declare that during the implementation of the research presented in the article, there was no conflict of interest.
Ethical Approval
This article does not contain any studies with human or animal participants performed by any of the authors.
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