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Surface modification of plasma nitriding on AlxCoCrFeNi high-entropy alloys
Journal of Materials Science & Technology 48 (2020) 140–145
Contents lists available at ScienceDirect
Journal of Materials Science & Technology journal homepage: www.jmst.org
Letter Surface modification of plasma nitriding on Alx CoCrFeNi high-entropy alloys
a r t i c l e Keywords: Plasma nitriding High-entropy alloy Microstructure Hardness Wear and tribology
i n f o
a b s t r a c t Plasma nitriding is successfully employed in treating Alx CoCrFeNi high-entropy alloys (HEAs) with finelydivided Al content (i.e., x values in molar ratio, x = 0.1- 0.8) to develop wear-resistant structural materials. Nitridation greatly removes the Al from the matrix that completely deplete the Ni-Al enriched phase, forming nanoscaled nitrides (AlN and CrN) precipitations near the surface. Nitriding promotes the hardness of present alloys with values widely ranging from 276 HV to 722 HV. Interestingly, the higher content the Al, the smaller thickness the nitrides layer, but the higher hardness due to the increased amount of hard nitrides phases and volume fraction of BCC phase. Significantly, plasma nitriding considerably improves the wear resistance of Alx CoCrFeNi HEAs by 4–18 times. © 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction Environmental issues are increasingly becoming a principal matter in the design of engines, nuclear, and tool materials. New materials, are therefore required in order to ensure the success of the novel industrial application. High-entropy alloys (HEAs) are newly designed and developed wherein the configurational entropy is higher than the fusion entropy of most conventional alloys due to an equal or near equal atomic composition with at least five components [1]. As one of the most thoroughly studied HEAs to date, i.e., Alx CoCrFeNi, presents a gradual structure transformation from face-centered-cubic (FCC) to body-centered-cubic (BCC) phase. The hardness of Alx CoCrFeNi HEAs is increased from 120 HV to 520 HV with raising Al content (x = 0–2.0) at as-cast states [2]. Commonly, plasma nitriding is applied mostly on steels [3], titanium alloys [4] and nickel alloys [5], and the surface mechanical performances are efficiently enhanced by the precipitated nitrides. Plasma nitriding therefore processes potential applicability in treating those HEAs containing large amounts of strong nitride forming elements such as Cr, Fe, Co, Ni, and Ti. Moreover, most of the nitirdes (AlN, Fe4 N, TiN, TaN, except for CrN) are refractory and show high thermal stability [5]. Particularly, the strongest binary nitride former, AlN, owning a lowest heat of formation binary nitrides (H = − 318 kJ/mol), compared with the other nitrides in this alloy system [6]. More importantly, with increasing the content of Al, the hardness of the nitrides layer is supposed to increase due to the increased amount of hard AlN [7]. In recent studies, plasma nitriding has been performed on HEAs comprising single phase, including CoCrFeMnNi (FCC) [8] and AlCoCrFeNi (BCC) HEAs [9]. Additionally, plasma nitriding was employed on promising dual-phase FeNiMnAlCr and Al1.3 CoCuFeNi2 HEAs (FCC + BCC) [10,11], Alx CoCrCuFeNi HEAs (x =
0–3.0) (two FCC) [7], and Al0.5 CrFe1.5 MnNi0.5 HEAs (two BCC) HEAs [12]. These dual-phase HEAs have significant potential in tool and structural industries on the consideration of their wide ranges of mechanical properties. As one of the mostly ambitious alloy systems, Alx CoCrFeNi HEAs present exceptional cryogenic toughness, excellent tensile strength, and well corrosion/wear resistance [13]. In order to further extending the application of this alloy system, plasma nitriding is eagerly utilized here to modify the surface performances of the current alloys. In the present work, plasma nitriding was conducted to improve the surface performances of Alx CoCrFeNi (x = 0.1 ∼ 0.8) alloys. The influences of Al content on the microstructural, mechanical, and wear properties of the nitrided HEAs were studied, and the corresponding mechanisms were explored in details. 2. Experimental One of the most widely studied system, i.e. Alx CoCrFeNi HEAs (x = 0.1, 0.25, 0.45, 0.6, 0.8) with different Al contents, were nitrided and investigated in this study. The samples were fabricated via vacuum arc melting and casting. The as-cast sheets were heated for homogenization treatment at 1100 ◦ C for 5 h in air atmosphere then quenched in water. The homogenized Al0.1, 0.25, 0.45 CoCrFeNi HEAs were cold rolled (CR) at room temperature to 70% thickness reductions, and the Al0.6,0.8 CoCrFeNi HEAs were cold rolled to 30% thickness reductions. All the as-rolled samples were recrystallized at 900 ◦ C for 1 h followed by water quenched. The recrystallized alloys were sliced into small specimens with a dimension of 10 mm × 5 mm, then ground, and well polished, using the 5 m grain size diamond-grinding paste to mirror. Subsequently, plasma nitriding was conducted in a DC glow discharge system operating at a constant voltage of 380 V at 550 ◦ C for 9 h. Meanwhile, the nitriding gas was ammonia with a flow rate
https://doi.org/10.1016/j.jmst.2020.01.057 1005-0302/© 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
Letter / Journal of Materials Science & Technology 48 (2020) 140–145
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Fig. 1. XRD patterns of (a) un-nitrided and (b) nitrided Alx CoCrFeNi HEAs.
of 0.4 m3 /h. The nitrides layer microstructures and thicknesses of the treated HEAs were investigated using a scanning electron microscope (SEM) equipped with an X-ray energy dispersive spectrometer (EDS). The crystal structures were characterized by using an X-ray diffractometer (XRD) with Cu K˛ radiation. Vickers microhardness tests were performed with an indentation load of 500 g holding for 15 s. The friction and wear tests were performed with a 5 mm Si3 N4 ball under dry sliding conditions by using a ball-on block high-speed reciprocating wear test apparatus (MFT-R4000). The normal load was 5 N and the sliding speed is 0.2 m/s. The reciprocating amplitude was set to 5 mm, and sliding time was fixed to 30 min. The wear loss was calculated using the depth and width of the wear track by a three-dimensional surface profiler based on scanning white light interferometry. 3. Results and discussion The detailed microstructural analysis were performed using XRD and SEM. The XRD patterns of the un-nitrided and nitrided Alx CoCrFeNi HEAs are shown in Fig. 1. As can be seen, only single FCC phase is detected for the Al0.1-0.25 CoCrFeNi HEAs, while both FCC and BCC phases are found in the Al0.45-0.8 CoCrFeNi HEAs. After nitriding, the intensity of BCC phase suddenly disappeared from the XRD patterns as soon as the nitrides is formed, which is consistent well with the early findings of nitrided Al1.3 CoCuFeNi2 and Al0.5 CrFe1.5 MnNi0.5 HEAs [11,14]. Similarly, nitridation greatly removes Al from the matrix of dual-phase FeNiMnAlCr HEAs (FCC + B2), resulting in a single-phase FCC matrix [10]. Note that all the nitrided HEAs indicate that the presence of some other phases. As the strongest nitride formers among the five elements, Al and Cr, own similar stoichiometric factor with nitrogen [5]. Here, the FCCstructured CrN and AlN are easily resolved via search-matching the JCPDS database whereas the HCP-structured AlN cannot be detected due to its smaller amount. Usually, CrN is a NaCl-type FCC structure whereas AlN is a wurtzite ZnS-type HCP structure after nitriding. Accordingly, the released nitrogen atoms diffuse into the Al alloys, forming AlN phase with a FCC structure in Ref. [15]. With
increasing x to 0.6-0.8, a minor of HCP-structured AlN phases begin to appear from XRD patterns. Among the five elements, nitrogen has a strong bonding energy with Al (−92 kJ/mol), which quickly leads to the formation of AlN. It must be mentioned that a ternary CrAlN coating, alloying CrN with varying Al content, has been successfully prepared by sputter etching [16]. Unambiguously, the XRD patterns of CrAlN coating (mixtures of CrN and AlN) are well consistent with present study. It is thus reasonable to express the two main nitrides as AlN and CrN, respectively. The microstructures of un-nitrided Alx CoCrFeNi HEAs display a phase transformation from a simple FCC solid-solution structure to a lamellar (FCC + B2) eutectic structure, as shown in Fig. 2. It can be easily seen that the Al0.1 CoCrFeNi and Al0.25 CoCrFeNi alloys have fully recrystallized microstructures with grain size are 15.2 ± 4.5 m and 12.0 ± 5.3 m after annealing at 900 ◦ C for 1 h, respectively. Annealing the Al0.45 CoCrFeNi alloy at 900 ◦ C results in a relatively fine-grained microstructure (1.5 ± 0.9 m) with many grain boundary precipitates, because the B2-type precipitates effectively retard the growth of grains during recrystallization [17]. The Al0.45 CoCrFeNi HEAs (Fig. 1(c)) are consist of approximately 90 vol-percent (vol.%) FCC phase and about 10 vol.% BCC/B2 phase. Duplex microstructures with a dendrite matrix (light gray regions) and a separated eutectic phase (gray regions) are presented in Al0.6-0.8 CoCrFeNi HEAs. The Al0.6 CoCrFeNi HEAs (Fig. 2(d)) are composed of the 64 vol.% dendritic FCC and 36 vol.% interdendritic BCC/B2 phases. The volume fraction of interdendritic BCC phase is increased to 48 vol.% with increasing the Al content to 0.8. The interdendrite is a Ni-Al enriched BCC phase, whereas the dendrite is a Cr enriched FCC phase identified from the EDS maps in Fig. 2(f) and (g). Note that the Al0.6 CoCrFeNi HEAs exhibit a relatively finer lamellar structure than the Al0.8 CoCrFeNi HEAs. Plenty of particle-shaped nitrides are uniformly distributed on the surface of nitrided Al0.1-0.45 CoCrFeNi HEAs with a variable size of (0.07-0.12) ± 0.03 m. As demonstrated in Fig. 3(d) and (e), the initial eutectic microstructures genetically produce the nitrided eutectic microstructures in Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs, respectively. Multiscaled particle-shaped nitrides are formed in the
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Letter / Journal of Materials Science & Technology 48 (2020) 140–145
Fig. 2. SEM images of the un-nitrided Alx CoCrFeNi HEAs: (a) x = 0.1; (b) x = 0.25; (c) x = 0.45; (d) x = 0.6; (e) x = 0.8. (f, g) Magnified microstructures of Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs with EDS elemental distribution maps, respectively.
Fig. 3. SEM images of nitrided Alx CoCrFeNi HEAs: (a) x = 0.1; (b) x = 0.25; (c) x = 0.45; (d) x = 0.6; (e) x = 0.8. (f, g) Magnified microstructures of Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs with EDS elemental distribution maps, respectively.
Letter / Journal of Materials Science & Technology 48 (2020) 140–145
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Fig. 4. SEM micrographs of the transverse cross-sections of nitrided Alx CoCrFeNi HEAs: (a) x = 0.1; (b) x = 0.25; (c) x = 0.45; (d) x = 0.6; (e) x = 0.8. (f) The relative chemical element intensity with the depth from the surface for Al0.1 CoCrFeNi HEAs.
eutectic Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs. The increased size of the nitrides at the interface between the dendrites and interdendrites compared to that both in dendrites and interdendrites is probably a consequence of the higher phase-boundary energy that increases the critical nucleus size and, hence, reduces the nucleation rate. The phase interface generates the large nitrides with size of 0.45 ± 0.11 m, attributing to the highly diffuse of nitrogen at the phase boundaries than that inside the crystals. The relatively finer nitrides particles with size of 0.10 ± 0.02 m are formed within the dendrites in Fig. 3g. In contrast, extremely small nitrides with size of 0.048 ± 0.012 m are precipitated in the interdendrite phase. The EDS analysis indicates that plasma nitriding balances the Al content in the Ni-Al-rich phase to that a FCC phase required. Fig. 4 shows the cross-sectional microstructures of plasma nitrided HEAs. Interestingly, with an increase in Al content, an obvious inhibition effect on nitrides depth with values varying widely from 15.5 ± 2.3 m to 5.5 ± 1.4 m occurs in the nitrided alloys. It can be clearly observed that the thickness of the nitrides layer is 15.5 ± 2.3 m for Al0.1 CoCrFeNi, which decreases to 5.5 ± 1.4 m for the Al0.8 CoCrFeNi HEAs. This is highly similar with the nitridation of Al-doped steels and HEAs: higher Al would reduce the thickness of nitrided layer [7,18]. As has been mentioned above, Al owns strong bonding with N, the much more Al content, the higher fraction of AlN in nitrides layer and thus reduces the effective cross-sectional region for N atom diffusion due to the topological effects. Hence, a higher amount of AlN would obstruct the diffusion path of N atoms and slow down the growth of nitrides layer. Furthermore, the larger Al atoms will distort the crystal lattice, which might also increases the energy barrier and thus the difficulty of nitrogen diffusion. As for the eutectic Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs, abundant of Ni-Al enriched BCC phases, having more open structure than the FCC structure, are observed before nitriding. But the results seem contradictory to the fact that the BCC phase has a higher interstitial diffusion rate than the FCC phase, which will generate
a larger nitrides layer. Interestingly, the diffusion of N along the phase boundaries results in a well nitrided layer with solid solution of nitrogen, as shown in Fig. 4(d) and (e). As have been previously demonstrated that the diffusion of N along phase boundaries is easier than that inside crystals (BCC and FCC phases), attributing to the higher phase-boundary energy. That’s to say, the lamellarstructured BCC and FCC phases will suppress the diffusion of N. The cross-sectional area for N atom diffusion in the relatively finer lamellar-structured Al0.6 CoCrFeNi HEAs is smaller than that in Al0.8 CoCrFeNi HEAs, i.e., the nitrided layer is reduced by the finer lamellar obstacles. The variation of chemical element concentration of nitrided Al0.1 CoCrFeNi HEAs with the depth from the surface is shown in Fig. 4(f). The penetration depth of the nitrogen into the substrate is in good agreement with the thickness of the nitrides layer. The results reported above indicate that the Al content has a strong influence on the nitridation rate in Alx CoCrFeNi HEAs. The microhardness of Alx CoCrFeNi HEAs as a function of Al content is displayed in Fig. 5. The hardness of the un-nitrided Al0.1 CoCrFeNi alloys is 188 HV, and it is increased to 320 HV with increasing Al content, attributing to the increased Ni-Al enriched BCC phases. Nitriding obviously improves the hardness of this alloy system, which is contributed to the hard nitrides and nitrided layer. Particularly, the Al0.8 CoCrFeNi HEAs present a high hardness of 722 HV after nitriding. The more Al content, the higher improved hardness, indicates the much larger content of AlN. It can be found that the thickness of nitrides area is decreased, as present in Fig. 4, but both the maximum surface hardening and substrate hardness are increased as the Al content increases. Generally, the wear rates of the metallic materials are evaluated according to Archard’s law [14], which is proportional to the hardness of the alloys. The ball-on-disk wear tests were curried out to examine wear resistances of the nitrided Alx CoCrFeNi HEAs. Fig. 6 shows the variation of the friction coefficient with sliding time for un-nitrided and nitride Alx CoCrFeNi HEAs. The friction coefficient
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Letter / Journal of Materials Science & Technology 48 (2020) 140–145 Table 1 Wear rates of unnitrided and nitrided Alx CoCrFeNi alloys.
Fig. 5. Variations of substrate hardness and BCC-volume fraction as a function of Al content in Alx CoCrFeNi alloys.
of the un-nitrided HEAs in air fluctuates around 0.21 and presents smaller fluctuations then becomes smooth until the end of the test. As shown in Fig. 6(a), that the friction coefficient of un-nitrided HEAs gradually ascends with increasing the Al content, attributing to the appearance of eutectic microstructures. Moreover, the curves of Al0.8 CoCrFeNi HEAs present larger fluctuation compared with the Al0.1 CoCrFeNi and Al0.45 CoCrFeNi HEAs due to a more serious periodically localized fracture of the eutectic surface layer accompanied with the periodic accumulation and elimination of the debris. Thus, some broad waves with relatively large fluctuation are observed with increasing the Al content in current alloy system. Significantly, plasma nitriding results in an increase in the friction coefficient for HEAs, as shown in Fig. 6(b). The precipitation of nanosacled nitrides improves the surface roughness of the alloys, which leads to a relatively higher friction coefficient. For the dual-phase HEAs, nitriding hardens both of FCC and BCC phases. The difference of the hardness between the two phases is greatly reduced. Hence, the periodic fluctuation of friction coefficient is obviously decreased. That’s to say, sliding friction now need to overcome the ultimate shear strength of the nitrides layer. Unambiguously, the ultimate shear strength of the nitrides layer is higher than the substrates of the un-nitride alloys. The tribological properties of nitrided HEAs therefore can
Alloy designation
Wear rate (mm3 /(N m))
Al0.1 CoCrFeNi Al0.25 CoCrFeNi Al0.45 CoCrFeNi Al0.6 CoCrFeNi Al0.8 CoCrFeNi
Unnitrided 1.61 × 10−4 1.47 × 10−4 2.54 × 10−4 2.53 × 10−4 2.51 × 10−4
Nitrided 3.95 × 10−5 1.72 × 10−5 1.65 × 10−5 1.54 × 10−5 1.33 × 10−5
Improved wear resistance (times) 4 9 15 16 18
be improved by precipitating the nano-nitrides, and the wear rate can be remarkably lowered, at room and high temperatures. Subsequently, the friction coefficient tends to be stable and exhibits periodic waves with slight fluctuation than the un-nitrided HEAs. The wear rate of the nitrided HEAs is lower than that of the un-nitrited HEAs under dry sliding, as listed in Table 1. For the single-FCC-phase Al0.1 CoCrFeNi and Al0.25 CoCrFeNi HEAs, the improved wear resistance is 4 and 9 times than the untreated ones, respectively. Surprisingly, the wear resistance of nitrided dualphase Al0.45-0.8 CoCrFeNi HEAs are highly enhanced by 15–18 times, compared with the un-nitrided ones. The structures and the functions of dry land ecosystem have been greatly changed under the synthetic effects of relatively large Al contents, phase components, and the amount of nitrides, resulting in a novelty phenomenon in this nitrided Alx CoCrFeNi HEAs. Accordingly, the harder nitrides layer formed by plasma nitriding leads to the reduction of the wear rate. What’s more, as the Al-varied coupling nitrides, series synergistic effects of improving the thermal stability, hardness, and damage resistance of CrN can be achieved [19,20]. 4. Conclusions In this study, Alx CoCrFeNi (x = 0.1-0.8) HEAs were plasma nitrided to study the influence of nitriding treatments on the microstructural and mechanical properties of the HEAs. The conclusions are summarized as follows: (1) The microstructure of Alx CoCrFeNi HEAs exhibit a gradual structure from single FCC to FCC and BCC phases. Nitriding
Fig. 6. Friction coefficient as a function of the sliding time for un-nitrided (a) and nitride (b) Alx CoCrFeNi HEAs.
Letter / Journal of Materials Science & Technology 48 (2020) 140–145
extremely reduces the intensity of BCC phase to zero. And AlN and CrN phases were formed on the surface of the nitride HEAs. (2) A well nitrides layer is observed here, which shows the thickness ranges from 15 to 5 m accompanied with hardness widely varying from 260 to 726 HV. With increasing the Al contents, the nitrides thicknesses is decreased due to the topological effect from AlN nitrides. (3) The average friction coefficient of current HEAs is increased from 0.21 to 0.37 with increasing Al contents from 0.1 to 0.8. In addition, nitriding increases the friction coefficient of HEAs in the dry condition due to the precipitated nitrides. (4) The wear tests revealed that the plasma nitriding process improves the wear resistance of Al0.1 CoCrFeNi HEAs considerably by 4–18 times. Acknowledgements This work was supported financially by the opening project from the National Key Laboratory for Remanufacturing (No. 61420050204) and Transformation of Scientific and Technological Achievements Programs of Higher Education Institutions in Shanxi (2019). J.X.H. would like to acknowledge the financial support from the China Scholarship Council (CSC). References [1] J.W. Yeh, S.K. Chen, S.J. Jin, J.Y. Gan, T.S. Chen, T.T. Shun, Adv. Eng. Mater. 6 (2004) 299–303. [2] W.R. Wang, W.L. Wang, S.C. Wang, Y.C. Tsai, C.H. Lai, J.W. Yeh, Intermetallics 26 (2012) 44–51. [3] B.S. Mann, V. Arya, Wear 249 (2001) 354–360. [4] Z.K. Song, H. Herman, Wear 80 (1982) 101–113. [5] U. Krupp, H.J. Christ, Oxid. Met. 52 (1999) 277–298. [6] C.T.L. Lynch, Handbook of Materials Science, CRC Press, Cleveland, 1974. [7] W.Y. Tang, J.W. Yeh, Metall. Mater. Trans. A 40 (2009) 1479–1486. [8] A. Nishimoto, T. Fukube, T. Maruyama, Surf. Coat.Technol. 376 (2019) 52–58. [9] Y. Wang, Y. Yang, H. Yang, M. Zhang, S. Ma, J. Qiao, Mater. Chem. Phys. 210 (2018) 233–239. [10] F. Meng, I. Baker, J. Alloys. Compd. 645 (2015) 376–381. [11] Y. Wang, Y. Yang, H. Yang, M. Zhang, J. Qiao, J. Alloys. Compd. 725 (2017) 365–372. [12] W.Y. Tang, M.H. Chuang, H.Y. Chen, J.W. Yeh, Surf. Coat. Technol. 204 (2010) 3118–3124.
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[13] D.B. Miracle, O.N. Senkov, Acta Mater. 122 (2017) 448–511. [14] Z. Tang, M.C. Gao, H. Diao, T. Yang, J. Liu, T. Zuo, Y. Zhang, Z. Lu, Y. Cheng, Y. Zhang, K.A. Dahmen, P.K. Liaw, T. Egami, JOM 65 (2013) 1848–1858. [15] K. Taherkhani, M. Soltanieh, J. Alloys. Compd. 741 (2018) 1247–1257. [16] Z. Li, P. Munroe, Z.T. Jiang, X. Zhao, J. Xu, Z.F. Zhou, J.Q. Jiang, F. Fang, Z.H. Xie, Acta Mater. 60 (2012) 5735–5744. [17] J. Hou, X. Shi, J. Qiao, Y. Zhang, P.K. Liaw, Y. Wu,Mater. Des. 180 (2019), 107910. [18] K.E. Thelning, Heat Treatment—Special, 6th ed., Butterworth-Heinemann, Boston, 1975, pp. 261–465. [19] J. Lin, J.J. Moore, B. Mishra, M. Pinkas, W.D. Sproul, J.A. Rees, Surf. Coat. Technol. 202 (2008) 1418–1436. [20] A.E. Reiter, C. Mitterer, B. Sartory, J. Vac, Sci. Technol. A 25 (2007) 711–720.
Jinxiong Hou College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China Wenwen Song Steel Institute, RWTH Aachen University, Aachen, 52072, Germany Liwei Lan College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China Junwei Qiao a,b,∗ College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China b Key Laboratory of Interface Science and Engineering in Advanced Materials, Ministry of Education, Taiyuan University of Technology, Taiyuan, 030024, China a
∗ Corresponding author at: College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China. E-mail address: [email protected] (J. Qiao)
28 September 2019 27 October 2019 31 January 2020 Available online 18 March 2020
Contents lists available at ScienceDirect
Journal of Materials Science & Technology journal homepage: www.jmst.org
Letter Surface modification of plasma nitriding on Alx CoCrFeNi high-entropy alloys
a r t i c l e Keywords: Plasma nitriding High-entropy alloy Microstructure Hardness Wear and tribology
i n f o
a b s t r a c t Plasma nitriding is successfully employed in treating Alx CoCrFeNi high-entropy alloys (HEAs) with finelydivided Al content (i.e., x values in molar ratio, x = 0.1- 0.8) to develop wear-resistant structural materials. Nitridation greatly removes the Al from the matrix that completely deplete the Ni-Al enriched phase, forming nanoscaled nitrides (AlN and CrN) precipitations near the surface. Nitriding promotes the hardness of present alloys with values widely ranging from 276 HV to 722 HV. Interestingly, the higher content the Al, the smaller thickness the nitrides layer, but the higher hardness due to the increased amount of hard nitrides phases and volume fraction of BCC phase. Significantly, plasma nitriding considerably improves the wear resistance of Alx CoCrFeNi HEAs by 4–18 times. © 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction Environmental issues are increasingly becoming a principal matter in the design of engines, nuclear, and tool materials. New materials, are therefore required in order to ensure the success of the novel industrial application. High-entropy alloys (HEAs) are newly designed and developed wherein the configurational entropy is higher than the fusion entropy of most conventional alloys due to an equal or near equal atomic composition with at least five components [1]. As one of the most thoroughly studied HEAs to date, i.e., Alx CoCrFeNi, presents a gradual structure transformation from face-centered-cubic (FCC) to body-centered-cubic (BCC) phase. The hardness of Alx CoCrFeNi HEAs is increased from 120 HV to 520 HV with raising Al content (x = 0–2.0) at as-cast states [2]. Commonly, plasma nitriding is applied mostly on steels [3], titanium alloys [4] and nickel alloys [5], and the surface mechanical performances are efficiently enhanced by the precipitated nitrides. Plasma nitriding therefore processes potential applicability in treating those HEAs containing large amounts of strong nitride forming elements such as Cr, Fe, Co, Ni, and Ti. Moreover, most of the nitirdes (AlN, Fe4 N, TiN, TaN, except for CrN) are refractory and show high thermal stability [5]. Particularly, the strongest binary nitride former, AlN, owning a lowest heat of formation binary nitrides (H = − 318 kJ/mol), compared with the other nitrides in this alloy system [6]. More importantly, with increasing the content of Al, the hardness of the nitrides layer is supposed to increase due to the increased amount of hard AlN [7]. In recent studies, plasma nitriding has been performed on HEAs comprising single phase, including CoCrFeMnNi (FCC) [8] and AlCoCrFeNi (BCC) HEAs [9]. Additionally, plasma nitriding was employed on promising dual-phase FeNiMnAlCr and Al1.3 CoCuFeNi2 HEAs (FCC + BCC) [10,11], Alx CoCrCuFeNi HEAs (x =
0–3.0) (two FCC) [7], and Al0.5 CrFe1.5 MnNi0.5 HEAs (two BCC) HEAs [12]. These dual-phase HEAs have significant potential in tool and structural industries on the consideration of their wide ranges of mechanical properties. As one of the mostly ambitious alloy systems, Alx CoCrFeNi HEAs present exceptional cryogenic toughness, excellent tensile strength, and well corrosion/wear resistance [13]. In order to further extending the application of this alloy system, plasma nitriding is eagerly utilized here to modify the surface performances of the current alloys. In the present work, plasma nitriding was conducted to improve the surface performances of Alx CoCrFeNi (x = 0.1 ∼ 0.8) alloys. The influences of Al content on the microstructural, mechanical, and wear properties of the nitrided HEAs were studied, and the corresponding mechanisms were explored in details. 2. Experimental One of the most widely studied system, i.e. Alx CoCrFeNi HEAs (x = 0.1, 0.25, 0.45, 0.6, 0.8) with different Al contents, were nitrided and investigated in this study. The samples were fabricated via vacuum arc melting and casting. The as-cast sheets were heated for homogenization treatment at 1100 ◦ C for 5 h in air atmosphere then quenched in water. The homogenized Al0.1, 0.25, 0.45 CoCrFeNi HEAs were cold rolled (CR) at room temperature to 70% thickness reductions, and the Al0.6,0.8 CoCrFeNi HEAs were cold rolled to 30% thickness reductions. All the as-rolled samples were recrystallized at 900 ◦ C for 1 h followed by water quenched. The recrystallized alloys were sliced into small specimens with a dimension of 10 mm × 5 mm, then ground, and well polished, using the 5 m grain size diamond-grinding paste to mirror. Subsequently, plasma nitriding was conducted in a DC glow discharge system operating at a constant voltage of 380 V at 550 ◦ C for 9 h. Meanwhile, the nitriding gas was ammonia with a flow rate
https://doi.org/10.1016/j.jmst.2020.01.057 1005-0302/© 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
Letter / Journal of Materials Science & Technology 48 (2020) 140–145
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Fig. 1. XRD patterns of (a) un-nitrided and (b) nitrided Alx CoCrFeNi HEAs.
of 0.4 m3 /h. The nitrides layer microstructures and thicknesses of the treated HEAs were investigated using a scanning electron microscope (SEM) equipped with an X-ray energy dispersive spectrometer (EDS). The crystal structures were characterized by using an X-ray diffractometer (XRD) with Cu K˛ radiation. Vickers microhardness tests were performed with an indentation load of 500 g holding for 15 s. The friction and wear tests were performed with a 5 mm Si3 N4 ball under dry sliding conditions by using a ball-on block high-speed reciprocating wear test apparatus (MFT-R4000). The normal load was 5 N and the sliding speed is 0.2 m/s. The reciprocating amplitude was set to 5 mm, and sliding time was fixed to 30 min. The wear loss was calculated using the depth and width of the wear track by a three-dimensional surface profiler based on scanning white light interferometry. 3. Results and discussion The detailed microstructural analysis were performed using XRD and SEM. The XRD patterns of the un-nitrided and nitrided Alx CoCrFeNi HEAs are shown in Fig. 1. As can be seen, only single FCC phase is detected for the Al0.1-0.25 CoCrFeNi HEAs, while both FCC and BCC phases are found in the Al0.45-0.8 CoCrFeNi HEAs. After nitriding, the intensity of BCC phase suddenly disappeared from the XRD patterns as soon as the nitrides is formed, which is consistent well with the early findings of nitrided Al1.3 CoCuFeNi2 and Al0.5 CrFe1.5 MnNi0.5 HEAs [11,14]. Similarly, nitridation greatly removes Al from the matrix of dual-phase FeNiMnAlCr HEAs (FCC + B2), resulting in a single-phase FCC matrix [10]. Note that all the nitrided HEAs indicate that the presence of some other phases. As the strongest nitride formers among the five elements, Al and Cr, own similar stoichiometric factor with nitrogen [5]. Here, the FCCstructured CrN and AlN are easily resolved via search-matching the JCPDS database whereas the HCP-structured AlN cannot be detected due to its smaller amount. Usually, CrN is a NaCl-type FCC structure whereas AlN is a wurtzite ZnS-type HCP structure after nitriding. Accordingly, the released nitrogen atoms diffuse into the Al alloys, forming AlN phase with a FCC structure in Ref. [15]. With
increasing x to 0.6-0.8, a minor of HCP-structured AlN phases begin to appear from XRD patterns. Among the five elements, nitrogen has a strong bonding energy with Al (−92 kJ/mol), which quickly leads to the formation of AlN. It must be mentioned that a ternary CrAlN coating, alloying CrN with varying Al content, has been successfully prepared by sputter etching [16]. Unambiguously, the XRD patterns of CrAlN coating (mixtures of CrN and AlN) are well consistent with present study. It is thus reasonable to express the two main nitrides as AlN and CrN, respectively. The microstructures of un-nitrided Alx CoCrFeNi HEAs display a phase transformation from a simple FCC solid-solution structure to a lamellar (FCC + B2) eutectic structure, as shown in Fig. 2. It can be easily seen that the Al0.1 CoCrFeNi and Al0.25 CoCrFeNi alloys have fully recrystallized microstructures with grain size are 15.2 ± 4.5 m and 12.0 ± 5.3 m after annealing at 900 ◦ C for 1 h, respectively. Annealing the Al0.45 CoCrFeNi alloy at 900 ◦ C results in a relatively fine-grained microstructure (1.5 ± 0.9 m) with many grain boundary precipitates, because the B2-type precipitates effectively retard the growth of grains during recrystallization [17]. The Al0.45 CoCrFeNi HEAs (Fig. 1(c)) are consist of approximately 90 vol-percent (vol.%) FCC phase and about 10 vol.% BCC/B2 phase. Duplex microstructures with a dendrite matrix (light gray regions) and a separated eutectic phase (gray regions) are presented in Al0.6-0.8 CoCrFeNi HEAs. The Al0.6 CoCrFeNi HEAs (Fig. 2(d)) are composed of the 64 vol.% dendritic FCC and 36 vol.% interdendritic BCC/B2 phases. The volume fraction of interdendritic BCC phase is increased to 48 vol.% with increasing the Al content to 0.8. The interdendrite is a Ni-Al enriched BCC phase, whereas the dendrite is a Cr enriched FCC phase identified from the EDS maps in Fig. 2(f) and (g). Note that the Al0.6 CoCrFeNi HEAs exhibit a relatively finer lamellar structure than the Al0.8 CoCrFeNi HEAs. Plenty of particle-shaped nitrides are uniformly distributed on the surface of nitrided Al0.1-0.45 CoCrFeNi HEAs with a variable size of (0.07-0.12) ± 0.03 m. As demonstrated in Fig. 3(d) and (e), the initial eutectic microstructures genetically produce the nitrided eutectic microstructures in Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs, respectively. Multiscaled particle-shaped nitrides are formed in the
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Fig. 2. SEM images of the un-nitrided Alx CoCrFeNi HEAs: (a) x = 0.1; (b) x = 0.25; (c) x = 0.45; (d) x = 0.6; (e) x = 0.8. (f, g) Magnified microstructures of Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs with EDS elemental distribution maps, respectively.
Fig. 3. SEM images of nitrided Alx CoCrFeNi HEAs: (a) x = 0.1; (b) x = 0.25; (c) x = 0.45; (d) x = 0.6; (e) x = 0.8. (f, g) Magnified microstructures of Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs with EDS elemental distribution maps, respectively.
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Fig. 4. SEM micrographs of the transverse cross-sections of nitrided Alx CoCrFeNi HEAs: (a) x = 0.1; (b) x = 0.25; (c) x = 0.45; (d) x = 0.6; (e) x = 0.8. (f) The relative chemical element intensity with the depth from the surface for Al0.1 CoCrFeNi HEAs.
eutectic Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs. The increased size of the nitrides at the interface between the dendrites and interdendrites compared to that both in dendrites and interdendrites is probably a consequence of the higher phase-boundary energy that increases the critical nucleus size and, hence, reduces the nucleation rate. The phase interface generates the large nitrides with size of 0.45 ± 0.11 m, attributing to the highly diffuse of nitrogen at the phase boundaries than that inside the crystals. The relatively finer nitrides particles with size of 0.10 ± 0.02 m are formed within the dendrites in Fig. 3g. In contrast, extremely small nitrides with size of 0.048 ± 0.012 m are precipitated in the interdendrite phase. The EDS analysis indicates that plasma nitriding balances the Al content in the Ni-Al-rich phase to that a FCC phase required. Fig. 4 shows the cross-sectional microstructures of plasma nitrided HEAs. Interestingly, with an increase in Al content, an obvious inhibition effect on nitrides depth with values varying widely from 15.5 ± 2.3 m to 5.5 ± 1.4 m occurs in the nitrided alloys. It can be clearly observed that the thickness of the nitrides layer is 15.5 ± 2.3 m for Al0.1 CoCrFeNi, which decreases to 5.5 ± 1.4 m for the Al0.8 CoCrFeNi HEAs. This is highly similar with the nitridation of Al-doped steels and HEAs: higher Al would reduce the thickness of nitrided layer [7,18]. As has been mentioned above, Al owns strong bonding with N, the much more Al content, the higher fraction of AlN in nitrides layer and thus reduces the effective cross-sectional region for N atom diffusion due to the topological effects. Hence, a higher amount of AlN would obstruct the diffusion path of N atoms and slow down the growth of nitrides layer. Furthermore, the larger Al atoms will distort the crystal lattice, which might also increases the energy barrier and thus the difficulty of nitrogen diffusion. As for the eutectic Al0.6 CoCrFeNi and Al0.8 CoCrFeNi HEAs, abundant of Ni-Al enriched BCC phases, having more open structure than the FCC structure, are observed before nitriding. But the results seem contradictory to the fact that the BCC phase has a higher interstitial diffusion rate than the FCC phase, which will generate
a larger nitrides layer. Interestingly, the diffusion of N along the phase boundaries results in a well nitrided layer with solid solution of nitrogen, as shown in Fig. 4(d) and (e). As have been previously demonstrated that the diffusion of N along phase boundaries is easier than that inside crystals (BCC and FCC phases), attributing to the higher phase-boundary energy. That’s to say, the lamellarstructured BCC and FCC phases will suppress the diffusion of N. The cross-sectional area for N atom diffusion in the relatively finer lamellar-structured Al0.6 CoCrFeNi HEAs is smaller than that in Al0.8 CoCrFeNi HEAs, i.e., the nitrided layer is reduced by the finer lamellar obstacles. The variation of chemical element concentration of nitrided Al0.1 CoCrFeNi HEAs with the depth from the surface is shown in Fig. 4(f). The penetration depth of the nitrogen into the substrate is in good agreement with the thickness of the nitrides layer. The results reported above indicate that the Al content has a strong influence on the nitridation rate in Alx CoCrFeNi HEAs. The microhardness of Alx CoCrFeNi HEAs as a function of Al content is displayed in Fig. 5. The hardness of the un-nitrided Al0.1 CoCrFeNi alloys is 188 HV, and it is increased to 320 HV with increasing Al content, attributing to the increased Ni-Al enriched BCC phases. Nitriding obviously improves the hardness of this alloy system, which is contributed to the hard nitrides and nitrided layer. Particularly, the Al0.8 CoCrFeNi HEAs present a high hardness of 722 HV after nitriding. The more Al content, the higher improved hardness, indicates the much larger content of AlN. It can be found that the thickness of nitrides area is decreased, as present in Fig. 4, but both the maximum surface hardening and substrate hardness are increased as the Al content increases. Generally, the wear rates of the metallic materials are evaluated according to Archard’s law [14], which is proportional to the hardness of the alloys. The ball-on-disk wear tests were curried out to examine wear resistances of the nitrided Alx CoCrFeNi HEAs. Fig. 6 shows the variation of the friction coefficient with sliding time for un-nitrided and nitride Alx CoCrFeNi HEAs. The friction coefficient
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Letter / Journal of Materials Science & Technology 48 (2020) 140–145 Table 1 Wear rates of unnitrided and nitrided Alx CoCrFeNi alloys.
Fig. 5. Variations of substrate hardness and BCC-volume fraction as a function of Al content in Alx CoCrFeNi alloys.
of the un-nitrided HEAs in air fluctuates around 0.21 and presents smaller fluctuations then becomes smooth until the end of the test. As shown in Fig. 6(a), that the friction coefficient of un-nitrided HEAs gradually ascends with increasing the Al content, attributing to the appearance of eutectic microstructures. Moreover, the curves of Al0.8 CoCrFeNi HEAs present larger fluctuation compared with the Al0.1 CoCrFeNi and Al0.45 CoCrFeNi HEAs due to a more serious periodically localized fracture of the eutectic surface layer accompanied with the periodic accumulation and elimination of the debris. Thus, some broad waves with relatively large fluctuation are observed with increasing the Al content in current alloy system. Significantly, plasma nitriding results in an increase in the friction coefficient for HEAs, as shown in Fig. 6(b). The precipitation of nanosacled nitrides improves the surface roughness of the alloys, which leads to a relatively higher friction coefficient. For the dual-phase HEAs, nitriding hardens both of FCC and BCC phases. The difference of the hardness between the two phases is greatly reduced. Hence, the periodic fluctuation of friction coefficient is obviously decreased. That’s to say, sliding friction now need to overcome the ultimate shear strength of the nitrides layer. Unambiguously, the ultimate shear strength of the nitrides layer is higher than the substrates of the un-nitride alloys. The tribological properties of nitrided HEAs therefore can
Alloy designation
Wear rate (mm3 /(N m))
Al0.1 CoCrFeNi Al0.25 CoCrFeNi Al0.45 CoCrFeNi Al0.6 CoCrFeNi Al0.8 CoCrFeNi
Unnitrided 1.61 × 10−4 1.47 × 10−4 2.54 × 10−4 2.53 × 10−4 2.51 × 10−4
Nitrided 3.95 × 10−5 1.72 × 10−5 1.65 × 10−5 1.54 × 10−5 1.33 × 10−5
Improved wear resistance (times) 4 9 15 16 18
be improved by precipitating the nano-nitrides, and the wear rate can be remarkably lowered, at room and high temperatures. Subsequently, the friction coefficient tends to be stable and exhibits periodic waves with slight fluctuation than the un-nitrided HEAs. The wear rate of the nitrided HEAs is lower than that of the un-nitrited HEAs under dry sliding, as listed in Table 1. For the single-FCC-phase Al0.1 CoCrFeNi and Al0.25 CoCrFeNi HEAs, the improved wear resistance is 4 and 9 times than the untreated ones, respectively. Surprisingly, the wear resistance of nitrided dualphase Al0.45-0.8 CoCrFeNi HEAs are highly enhanced by 15–18 times, compared with the un-nitrided ones. The structures and the functions of dry land ecosystem have been greatly changed under the synthetic effects of relatively large Al contents, phase components, and the amount of nitrides, resulting in a novelty phenomenon in this nitrided Alx CoCrFeNi HEAs. Accordingly, the harder nitrides layer formed by plasma nitriding leads to the reduction of the wear rate. What’s more, as the Al-varied coupling nitrides, series synergistic effects of improving the thermal stability, hardness, and damage resistance of CrN can be achieved [19,20]. 4. Conclusions In this study, Alx CoCrFeNi (x = 0.1-0.8) HEAs were plasma nitrided to study the influence of nitriding treatments on the microstructural and mechanical properties of the HEAs. The conclusions are summarized as follows: (1) The microstructure of Alx CoCrFeNi HEAs exhibit a gradual structure from single FCC to FCC and BCC phases. Nitriding
Fig. 6. Friction coefficient as a function of the sliding time for un-nitrided (a) and nitride (b) Alx CoCrFeNi HEAs.
Letter / Journal of Materials Science & Technology 48 (2020) 140–145
extremely reduces the intensity of BCC phase to zero. And AlN and CrN phases were formed on the surface of the nitride HEAs. (2) A well nitrides layer is observed here, which shows the thickness ranges from 15 to 5 m accompanied with hardness widely varying from 260 to 726 HV. With increasing the Al contents, the nitrides thicknesses is decreased due to the topological effect from AlN nitrides. (3) The average friction coefficient of current HEAs is increased from 0.21 to 0.37 with increasing Al contents from 0.1 to 0.8. In addition, nitriding increases the friction coefficient of HEAs in the dry condition due to the precipitated nitrides. (4) The wear tests revealed that the plasma nitriding process improves the wear resistance of Al0.1 CoCrFeNi HEAs considerably by 4–18 times. Acknowledgements This work was supported financially by the opening project from the National Key Laboratory for Remanufacturing (No. 61420050204) and Transformation of Scientific and Technological Achievements Programs of Higher Education Institutions in Shanxi (2019). J.X.H. would like to acknowledge the financial support from the China Scholarship Council (CSC). References [1] J.W. Yeh, S.K. Chen, S.J. Jin, J.Y. Gan, T.S. Chen, T.T. Shun, Adv. Eng. Mater. 6 (2004) 299–303. [2] W.R. Wang, W.L. Wang, S.C. Wang, Y.C. Tsai, C.H. Lai, J.W. Yeh, Intermetallics 26 (2012) 44–51. [3] B.S. Mann, V. Arya, Wear 249 (2001) 354–360. [4] Z.K. Song, H. Herman, Wear 80 (1982) 101–113. [5] U. Krupp, H.J. Christ, Oxid. Met. 52 (1999) 277–298. [6] C.T.L. Lynch, Handbook of Materials Science, CRC Press, Cleveland, 1974. [7] W.Y. Tang, J.W. Yeh, Metall. Mater. Trans. A 40 (2009) 1479–1486. [8] A. Nishimoto, T. Fukube, T. Maruyama, Surf. Coat.Technol. 376 (2019) 52–58. [9] Y. Wang, Y. Yang, H. Yang, M. Zhang, S. Ma, J. Qiao, Mater. Chem. Phys. 210 (2018) 233–239. [10] F. Meng, I. Baker, J. Alloys. Compd. 645 (2015) 376–381. [11] Y. Wang, Y. Yang, H. Yang, M. Zhang, J. Qiao, J. Alloys. Compd. 725 (2017) 365–372. [12] W.Y. Tang, M.H. Chuang, H.Y. Chen, J.W. Yeh, Surf. Coat. Technol. 204 (2010) 3118–3124.
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Jinxiong Hou College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China Wenwen Song Steel Institute, RWTH Aachen University, Aachen, 52072, Germany Liwei Lan College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China Junwei Qiao a,b,∗ College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China b Key Laboratory of Interface Science and Engineering in Advanced Materials, Ministry of Education, Taiyuan University of Technology, Taiyuan, 030024, China a
∗ Corresponding author at: College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China. E-mail address: [email protected] (J. Qiao)
28 September 2019 27 October 2019 31 January 2020 Available online 18 March 2020