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Thermally stable laser cladded CoCrCuFeNi high-entropy alloy coating with low stacking fault energy
Journal of Alloys and Compounds 600 (2014) 210–214
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
Thermally stable laser cladded CoCrCuFeNi high-entropy alloy coating with low stacking fault energy Hui Zhang a, Yi-Zhu He a, Ye Pan b, Sheng Guo c,⇑ a
School of Materials Science and Engineering, Anhui Key Laboratory of Materials Science and Processing, Anhui University of Technology, Ma’anshan 243002, Anhui, PR China School of Materials Science and Engineering, Jiangsu Key Laboratory of Advanced Metallic Materials, Southeast University, Nanjing 211189, Jiangsu, PR China c Surface and Microstructure Engineering Group, Department of Materials and Manufacturing Technology, Chalmers University of Technology, SE-41296, Gothenburg, Sweden b
a r t i c l e
i n f o
Article history: Received 29 January 2014 Received in revised form 19 February 2014 Accepted 20 February 2014 Available online 28 February 2014 Keywords: High-entropy alloys Laser cladding Thermal stability Stacking fault energy
a b s t r a c t The application of high-entropy alloys (HEAs) as coating materials has become an active research topic recently. Here an fcc structured CoCrCuFeNi HEA coating with a thickness of 1.2 mm was laser cladded onto a Q235 steel. The alloy coating possessed an excellent thermal stability in that no phase transformations occurred up to 1000 °C (0.86Tm), and the dendritic morphology of the as-solidified microstructure could be kept to higher than 750 °C (0.7Tm). After annealing the as-solidified coating at 750 °C for 5 h, the lattice distortion in the rapidly solidified alloy was reduced, resulting in a moderate decrease of both the hardness and electric resistivity. Interestingly, profuse stacking faults ribbons were observed in the dendritic region of the alloy after annealing, driven by the thermal stress. This phenomenon provided a direct experimental evidence of the low stacking fault energy in HEAs. The thermodynamic origin of the thermal stability for HEAs was proposed. Ó 2014 Elsevier B.V. All rights reserved.
1. Introduction The research on high-entropy alloys (HEAs) [1–3], or multicomponent alloys with equiatomic or close-to-equiatomic compositions, has become very active in recent years. On one hand, this new alloying strategy is a breakthrough to the conventional physical metallurgy, where using one or two principle elements has dominated the alloy development in the past. The massive opportunities created to develop new materials, to observe new phenomenon, and to explore new structural and functional properties are certainly attractive. On the other hand, the previous research on HEAs has shown great potentials of this new class of materials, in applications where high strength, high hardness, high wear resistance or high-temperature softening resistance are required [3–5]. Particularly, out of these potential applications of HEAs, using them as coating materials seems to be quite promising, from both technical and economic considerations. Indeed, there already exist many research activities in this respect. However, most of the previous coating preparations were using the magnetron sputtering method (e.g., see [6–10]), and the thickness of the achieved coating/film could reach only a couple of microns. Practically, methods like laser cladding [11–18] or
⇑ Corresponding author. Tel.: +46 (0)31 772 1254; fax: +46 (0)31 772 1313. E-mail address: [email protected] (S. Guo). http://dx.doi.org/10.1016/j.jallcom.2014.02.121 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.
thermal spray [19], which can prepare coatings with a thickness larger than 1 mm, are more desired if the target is the engineering applications. In this work, an fcc structured CoCrCuFeNi HEA coating with a thickness of 1.2 mm was prepared by the laser cladding method, to study its thermal stability and its potential to be used in the high-temperature environment. The target alloy composition has been studied extensively when it is prepared by the arc melting method (e.g., see [20–22]), which makes an easy comparison between laser cladded materials and arc-melted ones. The fcc structured alloy system was chosen because they have a densely packed structure and hence good creep properties [23]. Using fcc-structured HEAs for high-temperature coating materials has possibly another advantage: both experimental evidences [24–26] and theoretic calculations [27] have suggested that these materials have a low stacking fault energy (SFE), which is known to be beneficial for the creep properties at elevated temperatures. 2. Experimental The materials preparation procedure followed what was previously described in Ref. [12] and is briefly recapped here for clarity. A powder mixture of commercially pure (>99.7 wt.%) elemental powders of Co, Cr, Cu, Fe and Ni were mixed equiatomically by mechanically alloying. The particle sizes of the elemental powders ranged from 50 to 120 lm. To improve the coating quality and also to enhance the hardness, minor additives of Si (1.2 at.%), Mn (1.2 at.%) and Mo (2.8 at.%) were added when mixing the powders. The powder mixture was then placed onto the surface
H. Zhang et al. / Journal of Alloys and Compounds 600 (2014) 210–214
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of a Q235 steel to form a powder bed with the thickness of 1.7–2.0 mm. The pre-placed powder mixture was melted by the moving laser beam, and a coating with a thickness of 1.2 mm was formed onto the substrate. A 5 kW TJ-HLT5000 type continuous-wave CO2-laser system was used for cladding, with the high-purity argon supplied through the coaxial nozzle to prevent oxidation. The used laser power was 2.0 kW, the beam diameter was 4.5 mm, and the scanning speed was 400 mm/min. To investigate the thermal stability of the coating, the as-solidified alloy was annealed at 500, 750, and 1000 °C for 5 h, respectively. The phase constitution of the as-solidified and annealed alloys was identified by the X-ray diffratometer (XRD) using the Cu Ka radiation (Rigaku D/Max-rB). The microstructures were observed on finely polished and etched surfaces under the scanning electron microscope (SEM, JEOL JSM-6490), equipped with the energy dispersive spectrometer (EDS). The fine details of the microstructure were revealed by the transmission electron microscope (TEM, Tecnai G2, 200 kV). The TEM specimens of the coating were cut parallel to the substrate, and mechanically thinned then twin-jet electrochemically polished till perforation, using the 7 vol.% perchloric acid–93 vol.% ethanol solution. The Vickers hardness was measured on the polished surfaces by applying a load of 4.9 N for 20 s. The differential scanning calorimetric (DSC, Netzsch STA 449 F3) curve for the as-solidified alloy was measured from ambient temperature to 1350 °C, under the protection of a flow argon gas. The heating rate was 10 °C/min. The electrical resistivity of the as-solidified and annealed alloys was measured using a SZT-2 type four-probe instrument. The Q235 steel substrate was removed by grinding before the electrical resistivity measurement.
3. Results and discussion The XRD patterns for the as-solidified and 500, 750, 1000 °C annealed alloys are given in Fig. 1(a). Apparently, only one set of fcc typed diffraction peaks was observed in the as-solidified material, and the same phase constitution was kept even after annealing at 1000 °C for 5 h. Seen from the DSC curve as given in Fig. 1(b), this suggests that there existed no phase transformations for this alloy below the first melting phenomenon occurring at 1092 °C, which is believed to originate from the melting of the Cu-rich inter-dendritic phase with a lower melting point. The melting point of the alloy, assigned as the onset melting temperature of the second and also the last melting phenomenon [20], is 1201 °C or 1474 K. The XRD results then indicate that the laser-cladded coating has a thermal stability up to 0.86Tm (1000 °C). A closer look at the XRD diffraction peaks reveals that the peak position shifted to the right side after annealing, and this clearly says that the lattice constant reduced. The lattice constant for the as-solidified alloy, was 3.596 Å estimated from the position of the strongest diffraction peak (the (1 1 1) peak shown in the inset in Fig. 1(a)). It is noted here that the lattice constant here was larger than that of the arc-melting prepared alloys, 3.579 Å [20,22]. The expanding of the lattice constant shall very possibly come from the increased solid solutioning due to the faster cooling rate of the laser cladding method, compared to the arc-melting one. After annealing, the lattice strain due to the solid solutioning was released somehow by the atomic rearrangement (see discussion below on the composition analysis). The relaxation of the residual stress that was generated during the laser cladding process, with a faster cooling rate, could also contribute to this release of the lattice strain. For example, the lattice constant for the 750 °C/5 h annealed alloys was estimated to be 3.583 Å. The microstructure of the as-solidified and 750 °C/5 h annealed alloys are shown in Fig. 2. The as-solidified alloy has a typical dendritic structure, like many other cast HEAs. The compositional analysis for the dendritic phase and the inter-dendritic phase is given in Table 1. The inter-dendritic phase is enriched in Cu and depleted in Cr, compared to the dendritic phase, due to the segregation. The segregation was much less significant when comparing with that in the arc-melting prepared materials [20], owing to the faster cooling rate of the laser cladding method. The segregation apparently did not cause a significant change on the lattice constant, which accounted for only one set of fcc-typed peaks being detected by XRD. After annealing at 750 °C (0.7Tm) for 5 h, the morphology of the microstructure almost remained the same as that in the as-solidified alloy, and no obvious growth of
Fig. 1. (a) XRD patterns for as-solidified and 500, 750, 1000 °C/5 h annealed alloys, with the inset showing the enlarged diffraction peak around 2h = 40–47°; (b) DSC curve of the as-solidified alloy.
microstructural features could be observed. This again shows the excellent thermal stability of this laser-cladded alloy coating. The composition changed slightly after annealing, and particularly the enrichment of Cu and depletion of Cr was more obvious in the inter-dendritic phase. In addition, Ni was less seen in the inter-dendritic phase, and Fe was less seen in the dendritic phase. The compositional change certainly affected the atomic arrangement in the alloy, and is believed to lead to the lattice constant reduction. The atomic rearrangement very possibly released some of the lattice strain caused by the solid solutioning [22]. This argument is further supported by the hardness and electrical resistivity measurement shown in Fig. 3. Both the hardness and electrical resistivity reduced, indicating the reduced solid solutioning in annealed alloys. Certainly, the release of the residual stress during the annealing process could contribute to the reduced hardness. More specifically, the 5 h of annealing at 500 °C almost caused no change, a 5.5% reduction of hardness occurred after annealing at 750 °C for 5 h, and annealing at 1000 °C for 5 h led to a 16.3% reduction in hardness. Even after annealing at 1000 °C for 5 h, the hardness was still much higher than that of the arc-melting prepared alloy [28]. The mechanical stability, together with the morphological and phase stability for this laser cladded CoCrCuFeNi alloy, make it a good candidate coating material to be used at high temperatures, at least up to 750 °C. To further reveal the microstructural origin for the excellent thermal stability of this alloy, TEM studies were made on the 750 °C/5 h annealed specimen. Very interestingly, profuse ribbons of stacking faults (SFRs) were observed, as shown in Fig. 4(a) and (b). SFRs were only detected in the dendritic region (DR), and not in the inter-dendritic region (ID), where instead some spherical
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Fig. 3. Electrical resistivity and Vickers hardness for as-solidified and annealed alloys.
Fig. 2. Secondary electron images for (a) as-solidified and (b) 750 °C/5 h annealed alloys.
nano-precipitates were present (Fig. 4(b)). The selected-area electron diffraction patterns (SAED) taken from the ID region did not reveal any new phase other than the fcc phase (Fig. 4(c)). SAED for the SFRs (Fig. 4(d)) clearly showed streak-like features, characteristic for stacking faults. We believe that this is the first report of profuse stacking faults appearing in HEAs, and notably in the annealed but not mechanical deformed materials. The ease of stacking formation in this alloy provided direct experimental evidence to echo the theoretically predicted low stacking faults energy (SFE) in HEAs [27]. Previously, the formation of annealing twins [25,26] or deformation twins [24] was also claimed as evidences of the low SFE at HEAs. Twining, however, was not observed in the 750 °C/5 h annealed alloys. The appearance of stacking faults and at the same time the absence of twinning can possibly be explained by the formation mechanism of the stacking faults observed here. As mentioned above, SFRs were found only in the DR region, and not in the ID regions. Table 1 tells the information that the DR region has a higher concentration of Cr and lower concentration of Cu, compared to the ID region for the as-solidified alloy. Upon annealing at elevated temperatures, the much smaller thermal expansion coefficient of Cr (4.9 10 6/K) relative to Cu (16.5 10 6/K) [29] rendered the DR region in tension while the ID region in compression. The
compositions of the DR and ID region did change during the annealing process, but it only made the segregation more significant and would not change the thermal stress states in these regions. It is already known that the twin nucleation occurs only when certain conditions concerning the crystallographic orientation, sign of the applied stress and deformation level are satisfied [30]. The sign of the stress could possibly also explain why SFRs were only formed in the DR region which was under tension, and not in the ID region which was under compression. The failure to observer twins, on the other hand, could be due to that the level of thermal stress generated here was not significant to initiate the twinning. These assumptions apparently need more robust evidences to support, but they could be some starting points for the future work. There is another important issue in terms of the SFRs formation: they were not observed in the as-solidified condition, where the thermal stress also existed. The current authors tend to think that this is because the before mentioned atomic rearrangement was not involved in the as-solidified condition, and it is exactly the atomic rearrangement occurring during annealing that caused the formation of stacking faults. The faster cooling nature of the laser cladding process enabled a high level of residual stresses stored in the materials, the release of which through the atomic rearrangement actually increased the opportunity for the stacking faults to occur. This can possibly explain why SFRs were not previously seen in other cast and annealed HEAs. The formation of profuse stacking faults in this laser cladded CoCrCuFeNi alloy indicates the low SFE of this material. The lowering of SFE in alloys with increasing alloying concentration is normally accounted by the Suzuki interaction mechanism [31], basically arguing that it is due to the segregation of solute atoms to stacking faults. Recently, Yeh et al. suggested a new mechanism for the low SFE in HEAs [26]. They thought that in HEAs where there is no such a distinction between solvent and solutes, the segregation of so-called solute atoms is not required. They argued that the reduction of SFE in HEAs is not mainly due to the reduced
Table 1 EDS results for the as-solidified and 750 °C/5 h annealed CoCrCuFeNi alloy. Material
Elements (at.%) Co
Cr
Cu
Fe
Ni
Mn
Mo
Si
As-solidified Dendrite Inter-dendrite
19.86 18.45
23.32 13.54
13.78 23.85
19.77 20.11
18.27 19.05
1.04 1.19
2.79 2.61
1.17 1.20
750 °C/5 h Dendrite Inter-dendrite
18.34 19.25
29.10 13.34
12.06 26.85
16.43 19.65
19.09 15.04
1.11 1.67
2.34 2.86
1.53 1.34
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Fig. 4. (a) TEM image showing multiple stacking fault ribbons (SFRs) in the 750 °C/5 h annealing alloy. (b) Magnified TEM image further reviewing SFRs that are formed in the dendritic region (DR), not in the inter-dendritic region (ID). (c) and (d) are selected-area electron diffraction patterns taken from the ID and the SFR, respectively, as indicated in (a), respectively.
energy of stacking faults by solute segregation (the Suzuki mechanism), but mainly due to the increased energy of the distorted matrix, or the parts that are not affected by the stacking faults, as SFE is essentially the energy difference between these two parts (plus the interfacial energy [32]). Certainly, the compositional adjustment for the alloying atoms in the stacking faults can also reduce their energies, but the segregation is not required in this sense. The combined effect of the high energy of the distorted matrix and the reduced energy of the stacking faults lead to the low SFE in HEAs. The current authors are in general supporting this new mechanism, and we think the similar mechanism contributes to the thermal stability of the alloy. The driving force for the grain growth is the decrease of the grain boundary energy, and in this context it is the interfacial energy of the dendrites that controls the growth of dendrites. In HEAs where the whole-solute matrix [26] is severely distorted, and at the same time the rearrangement of the alloying elements at the interfacial regions releases the energy, the interfacial energy is reduced, in a similar way by which the SFE is reduced, giving a less driving force for the dendritic growth. Previously, the structural stability of HEAs is more attributed to the sluggish diffusion kinetics [33], and we now regard the thermodynamic factors also contributing to it. The alloys with a lower SFE are known to have better creep properties, because the more extended partial dislocations make crossslipping more difficult [34]. The above arguments all point to that HEAs with a low SFE are good candidate materials to be used in high-temperature environments. 4. Conclusions In summary, a fcc structured high-entropy CoCrCuFeNi alloy was laser cladded onto a Q235 steel substrate. Both the phase constitution and the dendritic morphology of the alloy coating showed
an impressive stability up to 0.7Tm (750 °C). Although the annealing induced relaxation of lattice strain caused a moderate decrease of hardness, a high level of hardness could be retained even after annealing at 1000 °C for 5 h. The thermal stress generated accompanying the annealing process resulted in the formation of profuse ribbons of stacking faults in the dendritic region, but no twins were observed. The low stacking fault energy accounts for the easy formation of stacking faults, and is also beneficial for the creep properties. The excellent thermal stability of high-entropy alloys has possibly a thermodynamic origin, in addition to the widely acknowledged kinetic one. Fcc-structured high-entropy alloys with a low stacking fault energy have the potential to be used as high-temperature coating materials. Acknowledgements The authors thank the financial support from the National Natural Science Foundation of China under Grant No. 51271001 and the Open Project of Jiangsu Key Laboratory of Advanced Metallic Materials under Grant No. AMM201201. SG thanks the initiation funding from the Area of Advance Materials Science at Chalmers University of Technology. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv. Eng. Mater. 6 (2004) 299–303. [2] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Mater. Sci. Eng. A 375–377 (2004) 213–218. [3] J.W. Yeh, Y.L. Chen, S.J. Lin, S.K. Chen, Mater. Sci. Forum 560 (2007) 1–9. [4] W.H. Wu, C.C. Yang, J.W. Yeh, Ann. Chim.-Sci. Mat. 31 (2006) 737–747. [5] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci. 61 (2014) 1–93. [6] T.K. Chen, T.T. Shun, J.W. Yeh, M.S. Wong, Surf. Coat. Technol. 188 (2004) 193– 200.
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[7] C.H. Lai, S.J. Lin, J.W. Yeh, S.Y. Chang, Surf. Coat. Technol. 201 (2006) 3275– 3280. [8] C.H. Lin, J.G. Duh, J.W. Yeh, Surf. Coat. Technol. 201 (2007) 6304–6308. [9] C.H. Lai, K.H. Cheng, S.J. Lin, J.W. Yeh, Surf. Coat. Technol. 202 (2008) 3732– 3738. [10] P.K. Huang, J.W. Yeh, Surf. Coat. Technol. 203 (2009) 1891–1896. [11] H. Zhang, Y. Pan, Y.Z. He, H.S. Jiao, Appl. Surf. Sci. 257 (2011) 2259–2263. [12] H. Zhang, Y. Pan, Y.Z. He, Mater. Des. 32 (2011) 1910–1915. [13] H. Zhang, Y. Pan, Y.Z. He, J. Therm. Spray Technol. 20 (2011) 1049–1055. [14] H. Zhang, Y.Z. He, Y. Pan, L.Z. Pei, Intermetallics 19 (2011) 1130–1135. [15] H. Zhang, Y. Pan, Y.Z. He, Surf. Coat. Technol. 205 (2011) 4068–4072. [16] Y.C. Lin, Y.H. Cho, Surf. Coat. Technol. 203 (2009) 1694–1701. [17] C. Huang, Y.Z. Zhang, J.Y. Shen, R. Vilar, Surf. Coat. Technol. 206 (2011) 1389– 1395. [18] C. Huang, Y.Z. Zhang, R. Vilar, J.Y. Shen, Mater. Des. 41 (2012) 338–343. [19] L.M. Wang, C.C. Chen, J.W. Yeh, S.T. Ke, Mater. Chem. Phys. 126 (2011) 880– 885. [20] C.J. Tong, Y.L. Chen, S.K. Chen, J.W. Yeh, T.T. Shun, C.H. Tsau, S.J. Lin, S.Y. Chang, Metall. Mater. Trans. A 36 (2005) 881–893.
[21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]
X.F. Wang, Y. Zhang, Y. Qiao, G.L. Chen, Intermetallics 15 (2007) 357–362. S. Guo, C. Ng, Z.J. Wang, C.T. Liu, J. Alloys Comp. 583 (2014) 410–413. F.R.N. Nabarro, H.L. De Villiers, The Physics of Creep, CRC Press, London, 1995. C.W. Tsai, Y.L. Chen, M.H. Tsai, J.W. Yeh, T.T. Shun, S.K. Chen, J. Alloys Comp. 486 (2009) 427–435. S.G. Ma, S.F. Zhang, M.C. Gao, P.K. Liaw, Y. Zhang, JOM 65 (2013) 1751– 1758. P.P. Bhattacharjee, G.D. Sathiaraj, M. Zaid, J.R. Gatti, C. Lee, C.W. Tsai, J.W. Yeh, J. Alloys Comp. 587 (2014) 544–552. A.J. Zaddach, C. Niu, C.C. Koch, D.L. Irving, JOM 65 (2013) 1780–1789. C.J. Tong, M.R. Chen, S.K. Chen, J.W. Yeh, T.T. Shun, S.J. Lin, S.Y. Chang, Metall. Mater. Trans. A 36 (2005) 1263–1271. http://www.webelements.com/. J.F.M. Vergnol, J.R. Grilhe, J. Phys. Fr. 45 (1984) 1479–1490. H. Suzuki, J. Phys. Soc. Jpn. 17 (1962) 322–325. W.S. Yang, C.M. Wan, J. Mater. Sci. 25 (1990) 1821–1823. K.Y. Tsai, M.H. Tsai, J.W. Yeh, Acta Mater. 61 (2013) 4887–4897. G.E. Dieter, Mechanical Metallurgy, McGraw-Hill, New York, 1976.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
Thermally stable laser cladded CoCrCuFeNi high-entropy alloy coating with low stacking fault energy Hui Zhang a, Yi-Zhu He a, Ye Pan b, Sheng Guo c,⇑ a
School of Materials Science and Engineering, Anhui Key Laboratory of Materials Science and Processing, Anhui University of Technology, Ma’anshan 243002, Anhui, PR China School of Materials Science and Engineering, Jiangsu Key Laboratory of Advanced Metallic Materials, Southeast University, Nanjing 211189, Jiangsu, PR China c Surface and Microstructure Engineering Group, Department of Materials and Manufacturing Technology, Chalmers University of Technology, SE-41296, Gothenburg, Sweden b
a r t i c l e
i n f o
Article history: Received 29 January 2014 Received in revised form 19 February 2014 Accepted 20 February 2014 Available online 28 February 2014 Keywords: High-entropy alloys Laser cladding Thermal stability Stacking fault energy
a b s t r a c t The application of high-entropy alloys (HEAs) as coating materials has become an active research topic recently. Here an fcc structured CoCrCuFeNi HEA coating with a thickness of 1.2 mm was laser cladded onto a Q235 steel. The alloy coating possessed an excellent thermal stability in that no phase transformations occurred up to 1000 °C (0.86Tm), and the dendritic morphology of the as-solidified microstructure could be kept to higher than 750 °C (0.7Tm). After annealing the as-solidified coating at 750 °C for 5 h, the lattice distortion in the rapidly solidified alloy was reduced, resulting in a moderate decrease of both the hardness and electric resistivity. Interestingly, profuse stacking faults ribbons were observed in the dendritic region of the alloy after annealing, driven by the thermal stress. This phenomenon provided a direct experimental evidence of the low stacking fault energy in HEAs. The thermodynamic origin of the thermal stability for HEAs was proposed. Ó 2014 Elsevier B.V. All rights reserved.
1. Introduction The research on high-entropy alloys (HEAs) [1–3], or multicomponent alloys with equiatomic or close-to-equiatomic compositions, has become very active in recent years. On one hand, this new alloying strategy is a breakthrough to the conventional physical metallurgy, where using one or two principle elements has dominated the alloy development in the past. The massive opportunities created to develop new materials, to observe new phenomenon, and to explore new structural and functional properties are certainly attractive. On the other hand, the previous research on HEAs has shown great potentials of this new class of materials, in applications where high strength, high hardness, high wear resistance or high-temperature softening resistance are required [3–5]. Particularly, out of these potential applications of HEAs, using them as coating materials seems to be quite promising, from both technical and economic considerations. Indeed, there already exist many research activities in this respect. However, most of the previous coating preparations were using the magnetron sputtering method (e.g., see [6–10]), and the thickness of the achieved coating/film could reach only a couple of microns. Practically, methods like laser cladding [11–18] or
⇑ Corresponding author. Tel.: +46 (0)31 772 1254; fax: +46 (0)31 772 1313. E-mail address: [email protected] (S. Guo). http://dx.doi.org/10.1016/j.jallcom.2014.02.121 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.
thermal spray [19], which can prepare coatings with a thickness larger than 1 mm, are more desired if the target is the engineering applications. In this work, an fcc structured CoCrCuFeNi HEA coating with a thickness of 1.2 mm was prepared by the laser cladding method, to study its thermal stability and its potential to be used in the high-temperature environment. The target alloy composition has been studied extensively when it is prepared by the arc melting method (e.g., see [20–22]), which makes an easy comparison between laser cladded materials and arc-melted ones. The fcc structured alloy system was chosen because they have a densely packed structure and hence good creep properties [23]. Using fcc-structured HEAs for high-temperature coating materials has possibly another advantage: both experimental evidences [24–26] and theoretic calculations [27] have suggested that these materials have a low stacking fault energy (SFE), which is known to be beneficial for the creep properties at elevated temperatures. 2. Experimental The materials preparation procedure followed what was previously described in Ref. [12] and is briefly recapped here for clarity. A powder mixture of commercially pure (>99.7 wt.%) elemental powders of Co, Cr, Cu, Fe and Ni were mixed equiatomically by mechanically alloying. The particle sizes of the elemental powders ranged from 50 to 120 lm. To improve the coating quality and also to enhance the hardness, minor additives of Si (1.2 at.%), Mn (1.2 at.%) and Mo (2.8 at.%) were added when mixing the powders. The powder mixture was then placed onto the surface
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of a Q235 steel to form a powder bed with the thickness of 1.7–2.0 mm. The pre-placed powder mixture was melted by the moving laser beam, and a coating with a thickness of 1.2 mm was formed onto the substrate. A 5 kW TJ-HLT5000 type continuous-wave CO2-laser system was used for cladding, with the high-purity argon supplied through the coaxial nozzle to prevent oxidation. The used laser power was 2.0 kW, the beam diameter was 4.5 mm, and the scanning speed was 400 mm/min. To investigate the thermal stability of the coating, the as-solidified alloy was annealed at 500, 750, and 1000 °C for 5 h, respectively. The phase constitution of the as-solidified and annealed alloys was identified by the X-ray diffratometer (XRD) using the Cu Ka radiation (Rigaku D/Max-rB). The microstructures were observed on finely polished and etched surfaces under the scanning electron microscope (SEM, JEOL JSM-6490), equipped with the energy dispersive spectrometer (EDS). The fine details of the microstructure were revealed by the transmission electron microscope (TEM, Tecnai G2, 200 kV). The TEM specimens of the coating were cut parallel to the substrate, and mechanically thinned then twin-jet electrochemically polished till perforation, using the 7 vol.% perchloric acid–93 vol.% ethanol solution. The Vickers hardness was measured on the polished surfaces by applying a load of 4.9 N for 20 s. The differential scanning calorimetric (DSC, Netzsch STA 449 F3) curve for the as-solidified alloy was measured from ambient temperature to 1350 °C, under the protection of a flow argon gas. The heating rate was 10 °C/min. The electrical resistivity of the as-solidified and annealed alloys was measured using a SZT-2 type four-probe instrument. The Q235 steel substrate was removed by grinding before the electrical resistivity measurement.
3. Results and discussion The XRD patterns for the as-solidified and 500, 750, 1000 °C annealed alloys are given in Fig. 1(a). Apparently, only one set of fcc typed diffraction peaks was observed in the as-solidified material, and the same phase constitution was kept even after annealing at 1000 °C for 5 h. Seen from the DSC curve as given in Fig. 1(b), this suggests that there existed no phase transformations for this alloy below the first melting phenomenon occurring at 1092 °C, which is believed to originate from the melting of the Cu-rich inter-dendritic phase with a lower melting point. The melting point of the alloy, assigned as the onset melting temperature of the second and also the last melting phenomenon [20], is 1201 °C or 1474 K. The XRD results then indicate that the laser-cladded coating has a thermal stability up to 0.86Tm (1000 °C). A closer look at the XRD diffraction peaks reveals that the peak position shifted to the right side after annealing, and this clearly says that the lattice constant reduced. The lattice constant for the as-solidified alloy, was 3.596 Å estimated from the position of the strongest diffraction peak (the (1 1 1) peak shown in the inset in Fig. 1(a)). It is noted here that the lattice constant here was larger than that of the arc-melting prepared alloys, 3.579 Å [20,22]. The expanding of the lattice constant shall very possibly come from the increased solid solutioning due to the faster cooling rate of the laser cladding method, compared to the arc-melting one. After annealing, the lattice strain due to the solid solutioning was released somehow by the atomic rearrangement (see discussion below on the composition analysis). The relaxation of the residual stress that was generated during the laser cladding process, with a faster cooling rate, could also contribute to this release of the lattice strain. For example, the lattice constant for the 750 °C/5 h annealed alloys was estimated to be 3.583 Å. The microstructure of the as-solidified and 750 °C/5 h annealed alloys are shown in Fig. 2. The as-solidified alloy has a typical dendritic structure, like many other cast HEAs. The compositional analysis for the dendritic phase and the inter-dendritic phase is given in Table 1. The inter-dendritic phase is enriched in Cu and depleted in Cr, compared to the dendritic phase, due to the segregation. The segregation was much less significant when comparing with that in the arc-melting prepared materials [20], owing to the faster cooling rate of the laser cladding method. The segregation apparently did not cause a significant change on the lattice constant, which accounted for only one set of fcc-typed peaks being detected by XRD. After annealing at 750 °C (0.7Tm) for 5 h, the morphology of the microstructure almost remained the same as that in the as-solidified alloy, and no obvious growth of
Fig. 1. (a) XRD patterns for as-solidified and 500, 750, 1000 °C/5 h annealed alloys, with the inset showing the enlarged diffraction peak around 2h = 40–47°; (b) DSC curve of the as-solidified alloy.
microstructural features could be observed. This again shows the excellent thermal stability of this laser-cladded alloy coating. The composition changed slightly after annealing, and particularly the enrichment of Cu and depletion of Cr was more obvious in the inter-dendritic phase. In addition, Ni was less seen in the inter-dendritic phase, and Fe was less seen in the dendritic phase. The compositional change certainly affected the atomic arrangement in the alloy, and is believed to lead to the lattice constant reduction. The atomic rearrangement very possibly released some of the lattice strain caused by the solid solutioning [22]. This argument is further supported by the hardness and electrical resistivity measurement shown in Fig. 3. Both the hardness and electrical resistivity reduced, indicating the reduced solid solutioning in annealed alloys. Certainly, the release of the residual stress during the annealing process could contribute to the reduced hardness. More specifically, the 5 h of annealing at 500 °C almost caused no change, a 5.5% reduction of hardness occurred after annealing at 750 °C for 5 h, and annealing at 1000 °C for 5 h led to a 16.3% reduction in hardness. Even after annealing at 1000 °C for 5 h, the hardness was still much higher than that of the arc-melting prepared alloy [28]. The mechanical stability, together with the morphological and phase stability for this laser cladded CoCrCuFeNi alloy, make it a good candidate coating material to be used at high temperatures, at least up to 750 °C. To further reveal the microstructural origin for the excellent thermal stability of this alloy, TEM studies were made on the 750 °C/5 h annealed specimen. Very interestingly, profuse ribbons of stacking faults (SFRs) were observed, as shown in Fig. 4(a) and (b). SFRs were only detected in the dendritic region (DR), and not in the inter-dendritic region (ID), where instead some spherical
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Fig. 3. Electrical resistivity and Vickers hardness for as-solidified and annealed alloys.
Fig. 2. Secondary electron images for (a) as-solidified and (b) 750 °C/5 h annealed alloys.
nano-precipitates were present (Fig. 4(b)). The selected-area electron diffraction patterns (SAED) taken from the ID region did not reveal any new phase other than the fcc phase (Fig. 4(c)). SAED for the SFRs (Fig. 4(d)) clearly showed streak-like features, characteristic for stacking faults. We believe that this is the first report of profuse stacking faults appearing in HEAs, and notably in the annealed but not mechanical deformed materials. The ease of stacking formation in this alloy provided direct experimental evidence to echo the theoretically predicted low stacking faults energy (SFE) in HEAs [27]. Previously, the formation of annealing twins [25,26] or deformation twins [24] was also claimed as evidences of the low SFE at HEAs. Twining, however, was not observed in the 750 °C/5 h annealed alloys. The appearance of stacking faults and at the same time the absence of twinning can possibly be explained by the formation mechanism of the stacking faults observed here. As mentioned above, SFRs were found only in the DR region, and not in the ID regions. Table 1 tells the information that the DR region has a higher concentration of Cr and lower concentration of Cu, compared to the ID region for the as-solidified alloy. Upon annealing at elevated temperatures, the much smaller thermal expansion coefficient of Cr (4.9 10 6/K) relative to Cu (16.5 10 6/K) [29] rendered the DR region in tension while the ID region in compression. The
compositions of the DR and ID region did change during the annealing process, but it only made the segregation more significant and would not change the thermal stress states in these regions. It is already known that the twin nucleation occurs only when certain conditions concerning the crystallographic orientation, sign of the applied stress and deformation level are satisfied [30]. The sign of the stress could possibly also explain why SFRs were only formed in the DR region which was under tension, and not in the ID region which was under compression. The failure to observer twins, on the other hand, could be due to that the level of thermal stress generated here was not significant to initiate the twinning. These assumptions apparently need more robust evidences to support, but they could be some starting points for the future work. There is another important issue in terms of the SFRs formation: they were not observed in the as-solidified condition, where the thermal stress also existed. The current authors tend to think that this is because the before mentioned atomic rearrangement was not involved in the as-solidified condition, and it is exactly the atomic rearrangement occurring during annealing that caused the formation of stacking faults. The faster cooling nature of the laser cladding process enabled a high level of residual stresses stored in the materials, the release of which through the atomic rearrangement actually increased the opportunity for the stacking faults to occur. This can possibly explain why SFRs were not previously seen in other cast and annealed HEAs. The formation of profuse stacking faults in this laser cladded CoCrCuFeNi alloy indicates the low SFE of this material. The lowering of SFE in alloys with increasing alloying concentration is normally accounted by the Suzuki interaction mechanism [31], basically arguing that it is due to the segregation of solute atoms to stacking faults. Recently, Yeh et al. suggested a new mechanism for the low SFE in HEAs [26]. They thought that in HEAs where there is no such a distinction between solvent and solutes, the segregation of so-called solute atoms is not required. They argued that the reduction of SFE in HEAs is not mainly due to the reduced
Table 1 EDS results for the as-solidified and 750 °C/5 h annealed CoCrCuFeNi alloy. Material
Elements (at.%) Co
Cr
Cu
Fe
Ni
Mn
Mo
Si
As-solidified Dendrite Inter-dendrite
19.86 18.45
23.32 13.54
13.78 23.85
19.77 20.11
18.27 19.05
1.04 1.19
2.79 2.61
1.17 1.20
750 °C/5 h Dendrite Inter-dendrite
18.34 19.25
29.10 13.34
12.06 26.85
16.43 19.65
19.09 15.04
1.11 1.67
2.34 2.86
1.53 1.34
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Fig. 4. (a) TEM image showing multiple stacking fault ribbons (SFRs) in the 750 °C/5 h annealing alloy. (b) Magnified TEM image further reviewing SFRs that are formed in the dendritic region (DR), not in the inter-dendritic region (ID). (c) and (d) are selected-area electron diffraction patterns taken from the ID and the SFR, respectively, as indicated in (a), respectively.
energy of stacking faults by solute segregation (the Suzuki mechanism), but mainly due to the increased energy of the distorted matrix, or the parts that are not affected by the stacking faults, as SFE is essentially the energy difference between these two parts (plus the interfacial energy [32]). Certainly, the compositional adjustment for the alloying atoms in the stacking faults can also reduce their energies, but the segregation is not required in this sense. The combined effect of the high energy of the distorted matrix and the reduced energy of the stacking faults lead to the low SFE in HEAs. The current authors are in general supporting this new mechanism, and we think the similar mechanism contributes to the thermal stability of the alloy. The driving force for the grain growth is the decrease of the grain boundary energy, and in this context it is the interfacial energy of the dendrites that controls the growth of dendrites. In HEAs where the whole-solute matrix [26] is severely distorted, and at the same time the rearrangement of the alloying elements at the interfacial regions releases the energy, the interfacial energy is reduced, in a similar way by which the SFE is reduced, giving a less driving force for the dendritic growth. Previously, the structural stability of HEAs is more attributed to the sluggish diffusion kinetics [33], and we now regard the thermodynamic factors also contributing to it. The alloys with a lower SFE are known to have better creep properties, because the more extended partial dislocations make crossslipping more difficult [34]. The above arguments all point to that HEAs with a low SFE are good candidate materials to be used in high-temperature environments. 4. Conclusions In summary, a fcc structured high-entropy CoCrCuFeNi alloy was laser cladded onto a Q235 steel substrate. Both the phase constitution and the dendritic morphology of the alloy coating showed
an impressive stability up to 0.7Tm (750 °C). Although the annealing induced relaxation of lattice strain caused a moderate decrease of hardness, a high level of hardness could be retained even after annealing at 1000 °C for 5 h. The thermal stress generated accompanying the annealing process resulted in the formation of profuse ribbons of stacking faults in the dendritic region, but no twins were observed. The low stacking fault energy accounts for the easy formation of stacking faults, and is also beneficial for the creep properties. The excellent thermal stability of high-entropy alloys has possibly a thermodynamic origin, in addition to the widely acknowledged kinetic one. Fcc-structured high-entropy alloys with a low stacking fault energy have the potential to be used as high-temperature coating materials. Acknowledgements The authors thank the financial support from the National Natural Science Foundation of China under Grant No. 51271001 and the Open Project of Jiangsu Key Laboratory of Advanced Metallic Materials under Grant No. AMM201201. SG thanks the initiation funding from the Area of Advance Materials Science at Chalmers University of Technology. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv. Eng. Mater. 6 (2004) 299–303. [2] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Mater. Sci. Eng. A 375–377 (2004) 213–218. [3] J.W. Yeh, Y.L. Chen, S.J. Lin, S.K. Chen, Mater. Sci. Forum 560 (2007) 1–9. [4] W.H. Wu, C.C. Yang, J.W. Yeh, Ann. Chim.-Sci. Mat. 31 (2006) 737–747. [5] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci. 61 (2014) 1–93. [6] T.K. Chen, T.T. Shun, J.W. Yeh, M.S. Wong, Surf. Coat. Technol. 188 (2004) 193– 200.
214
H. Zhang et al. / Journal of Alloys and Compounds 600 (2014) 210–214
[7] C.H. Lai, S.J. Lin, J.W. Yeh, S.Y. Chang, Surf. Coat. Technol. 201 (2006) 3275– 3280. [8] C.H. Lin, J.G. Duh, J.W. Yeh, Surf. Coat. Technol. 201 (2007) 6304–6308. [9] C.H. Lai, K.H. Cheng, S.J. Lin, J.W. Yeh, Surf. Coat. Technol. 202 (2008) 3732– 3738. [10] P.K. Huang, J.W. Yeh, Surf. Coat. Technol. 203 (2009) 1891–1896. [11] H. Zhang, Y. Pan, Y.Z. He, H.S. Jiao, Appl. Surf. Sci. 257 (2011) 2259–2263. [12] H. Zhang, Y. Pan, Y.Z. He, Mater. Des. 32 (2011) 1910–1915. [13] H. Zhang, Y. Pan, Y.Z. He, J. Therm. Spray Technol. 20 (2011) 1049–1055. [14] H. Zhang, Y.Z. He, Y. Pan, L.Z. Pei, Intermetallics 19 (2011) 1130–1135. [15] H. Zhang, Y. Pan, Y.Z. He, Surf. Coat. Technol. 205 (2011) 4068–4072. [16] Y.C. Lin, Y.H. Cho, Surf. Coat. Technol. 203 (2009) 1694–1701. [17] C. Huang, Y.Z. Zhang, J.Y. Shen, R. Vilar, Surf. Coat. Technol. 206 (2011) 1389– 1395. [18] C. Huang, Y.Z. Zhang, R. Vilar, J.Y. Shen, Mater. Des. 41 (2012) 338–343. [19] L.M. Wang, C.C. Chen, J.W. Yeh, S.T. Ke, Mater. Chem. Phys. 126 (2011) 880– 885. [20] C.J. Tong, Y.L. Chen, S.K. Chen, J.W. Yeh, T.T. Shun, C.H. Tsau, S.J. Lin, S.Y. Chang, Metall. Mater. Trans. A 36 (2005) 881–893.
[21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]
X.F. Wang, Y. Zhang, Y. Qiao, G.L. Chen, Intermetallics 15 (2007) 357–362. S. Guo, C. Ng, Z.J. Wang, C.T. Liu, J. Alloys Comp. 583 (2014) 410–413. F.R.N. Nabarro, H.L. De Villiers, The Physics of Creep, CRC Press, London, 1995. C.W. Tsai, Y.L. Chen, M.H. Tsai, J.W. Yeh, T.T. Shun, S.K. Chen, J. Alloys Comp. 486 (2009) 427–435. S.G. Ma, S.F. Zhang, M.C. Gao, P.K. Liaw, Y. Zhang, JOM 65 (2013) 1751– 1758. P.P. Bhattacharjee, G.D. Sathiaraj, M. Zaid, J.R. Gatti, C. Lee, C.W. Tsai, J.W. Yeh, J. Alloys Comp. 587 (2014) 544–552. A.J. Zaddach, C. Niu, C.C. Koch, D.L. Irving, JOM 65 (2013) 1780–1789. C.J. Tong, M.R. Chen, S.K. Chen, J.W. Yeh, T.T. Shun, S.J. Lin, S.Y. Chang, Metall. Mater. Trans. A 36 (2005) 1263–1271. http://www.webelements.com/. J.F.M. Vergnol, J.R. Grilhe, J. Phys. Fr. 45 (1984) 1479–1490. H. Suzuki, J. Phys. Soc. Jpn. 17 (1962) 322–325. W.S. Yang, C.M. Wan, J. Mater. Sci. 25 (1990) 1821–1823. K.Y. Tsai, M.H. Tsai, J.W. Yeh, Acta Mater. 61 (2013) 4887–4897. G.E. Dieter, Mechanical Metallurgy, McGraw-Hill, New York, 1976.