Effect of high-temperature exposure on the microstructure and mechanical properties of the Al5Ti5Co35Ni35Fe20 high-entropy alloy

j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 www.jmrt.com.br Available online at www.sciencedirect.com Original Article Effect of high-temperature exposure on the microstructure and mechanical properties of the Al5Ti5Co35Ni35Fe20 high-entropy alloy Piotr Bałaa,b,∗, Kamil Góreckic, Wiktor Bednarczyka, Maria W˛ atrobaa, Sebastian Lecha, Jakub Kawałkob aAGH University of Science and Technology, Faculty of Metals Engineering and Industrial Computer Science, Al. A. Mickiewicza 30, 30-059 Krakow, Poland bAGH University of Science and Technology, Academic Centre for Materials and Nanotechnology, Al. A. Mickiewicza 30, 30-059 Krakow, Poland cVSB - Technical University of Ostrava, Institute of Environmental Technology, 17. listopadu 2172/15, 708 00 Ostrava-Poruba, Ostrava, Czech Republic a r t i c l e i n f o Article history: Received 23 July 2019 Accepted 31 October 2019 Available online 23 November 2019 Keywords: High entropy alloys Aging Microstructure ␥ Discontinuous precipitation a b s t r a c t Effect of high-temperature exposure on the microstructure and mechanical properties of the Al5Ti5Co35Ni35Fe20 High-Entropy Alloy was studied. High-entropy alloys belong to the group of multi-principal element alloys composed of at least five elements with a concentration between 5 and 35 atomic pct. Despite the multicomponent composition, the goal is to achieve a simple solid solution. The alloy was exposed to a temperature range of 650–900 ◦C for 168 h (7 days), after hot rolling and air-quenching. After hot rolling the investigated alloy was not fully supersaturated. The application of air cooling caused ␥precipitation, which was uniformly distributed within the microstructure. High-temperature exposure caused the growth of ␥particles. At 700 ◦C and higher, additional continuous and discontinuous precipitation of ␥was observed. Finally, the microstructure of the investigated alloy consisted of ␥ matrix and three types of ␥’ precipitates, i.e. spherical ␥phases with bimodal distribution within the grains and elongated DP ␥phases present at the grain boundaries. The highest mechanical properties were obtained after exposure to 650 ◦C, which was caused by ␥precipitation. Further increase in temperature decreased the mechanical properties. © 2019 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/). ∗Corresponding author. E-mail: [email protected] (P. Bała). 1. Introduction Metallic materials are frequently chosen for structural applications as they have a desirable combination of mechanical characteristics. The final properties of alloys are often the result of not one, but two or more strengthening mechanisms. In this group there are many alloys like stainless steels, https://doi.org/10.1016/j.jmrt.2019.10.084 2238-7854/© 2019 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http:// creativecommons.org/licenses/by-nc-nd/4.0/). 552 j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 high-speed steels, superalloys, and multi-principal element alloys. Usually, they have a complex chemical composition and require advanced production technologies [1–4]. High-Entropy Alloys (HEAs) belong to the group of multiprincipal element alloys that are composed of at least five elements with the concentration of each element between 5 and 35 atomic pct. This new group of materials is characterized by several so-called core effects like the high-entropy effect, severe lattice distortion effect, sluggish diffusion effect and cocktail effect. These core effects are responsible for unique properties, including high hardness, thermal resistance, abrasion resistance, and others. One of the most interesting features, resulting directly from the chemical composition of HEAs, is that despite the multicomponent composition, only simple phases form. The goal is to achieve a simple solid solution like face-centered cubic (FCC) or body-centered cubic (BCC) or both (FCC + BCC) [5–7]. It has been reported that in HEA’s, solid solutions form rather than intermetallic compounds or other complex ordered phases owing to the high entropy of mixing. The high entropy of mixing stabilized solid solution is very strong, and it is not possible to indicate, which element is the solvent and which ones are solutes [5,6]. HEAs could be in an equilibrium state if sufficiently annealed. However, this is generally a lengthy process in comparison to conventional alloys, due to lower diffusion rates of alloying elements and phase transformation rates [8]. This makes the synthesis of stable HEAs, consisting only of a single phase solid solution, nearly impossible, even in the case of the Cantor alloy, in which authors [9–11] observed ␴ phase precipitation. However, some authors [12] suggest, that new materials, which are not fully characterized as solutions, should not be disqualified. Often, materials with a high-entropy matrix and additional precipitates possess huge application potential. Therefore, HEAs could have a broad range of chemical compositions, phases, and microstructures and thus generate the desired physical, mechanical, and chemical properties. Certainly, controlling metastable structures to obtain the best performance for specific applications is an important research issue of HEAs [13]. The goal of our research is to investigate the microstructural changes during long term high-temperature exposure of HEA’s, which contains elements such as Al or Ti, prone to create intermetallic phases instead of a solid solution. This knowledge is crucial when considering a material for hightemperature applications. 2. Materials and methods The alloy, with the nominal chemical composition of (at. pct.): Al–5 pct., Ti–5 pct., Co–35 pct., Ni–35 pct., Fe–20 pct., was synthesized in a vacuum melting furnace. Next, a rodshaped ingot was heat treated in a vacuum furnace for 20 h at 1100 ◦C to obtain chemical homogeneity throughout the entire volume. Slabs were supersaturated for 1 h at 1200 ◦C and subsequently water quenched. After the supersaturation process, the slabs (4 mm of thickness) were hot rolled (heating to 1100 ◦C after every rolling pass) for thickness reduction to 1 mm. After the final pass, the material was quenched in air. Table 1 – Thermodynamic factors and VEC parameter for the investigated alloy. Hmix, kJ/mol ı, % ˝, − VEC,− −9.5 4.8 2.0 8.5 Finally, the material was exposure to a temperature range of 650–900 ◦C for168 h (7 days). Tensile flat dog-bone specimens were cut using a wire electric discharge machine with a gauge dimension of 15 × 3 × 1 mm. Tensile tests were conducted via a universal testing machine (INSTRON 5966) at a strain rate of 2 × 10−3 s−1at room temperature. Microstructural investigations were carried out by means of a light microscope (LM–Nikon Eclipse LV150 N), scanning electron microscope (SEM–FEI Versa 3D) with an EDAX electron backscatter diffraction (EBSD) system, and a transmission electron microscope (TEM - FEI Tecnai TF-20 X-TWIN). An EBSD map was collected from a 145 ␮m × 125 ␮m area of the hot-rolled sample with a step size of 100 nm. Samples for LM and SEM observations were prepared using standard metallographic techniques and polished using silica oxides with a small addition of malic acid. Next, the samples were etched using a 4 g CuSO4+ 20 ml HCl + 20 ml C2H5OH solution for about 2 s. Thin foils for TEM investigations were mechanically prepared and electropolished in a 10-pct. solution of perchloric acid in methanol, at −12 ◦C and 20 V. It is worth mentioning that the investigated alloy is strongly ferromagnetic, which caused many problems, especially during TEM analysis. It is well known that the Ni-Fe-Co alloys have good ferromagnetic properties [15,18]. The crystallographic structure of the hot rolled material was analyzed by mean of X-ray diffraction with CoK␣ radiation (Panalytical Empyrean). Hardness measurements were performed by means of an automated hardness tester (Tukon 2500 Willson-Hardness) and Vickers indenter with the applied force of 9.81 N. 3. Results and discussion Yang et al. presented in [16] the thermodynamic factors required for high-entropy alloys, namely mixing enthalpy (Hmix between -15 and -5 kJ/mol), atomic ratio difference parameter ı (≤6.6%), and enthalpy to entropy ratio ˝ (≥1.1). In our previous paper [17], we calculated the factors (see Table 1) and our alloy fulfills the above-mentioned requirements. Additionally, the Valance Electron Concentration (VEC) proposed by Guo in [14] (when VEC value: lower than 6.8–BCC structure, higher than 8–FCC structure, and between BCC + FCC structure) was calculated, suggesting that the alloy should have a face-centered cubic structure. The XRD pattern of the investigated alloy after rolling and air quenching is presented in Fig. 1. Only peaks related to FCC solid solution in position 51, 59 and 89◦2 were recorded, which is in agreement with the theoretical prediction (VEC). The microstructure of the investigated alloy after rolling is presented in Fig. 2. The grain size in this state varies between 50 and 100 ␮m. Numerous aging twins (see Fig. 2b-d) and subgrain structures in the form of intragranular misorientations (see. Fig. 2b) observed in the microstructure seem to correspond with the data presented by Gao et al. [8], related to j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 553 Fig. 1 – XRD pattern of the investigated alloy in as-hot rolled state and after high-temperature exposure to the temperature range of 650–900 ◦C for 168 h. problems with achieving equilibrium state in HEA’s. Additionally, during air cooling applied after hot rolling, the ␥(A3B phase) precipitation occurred. The fine, not exceeding 15 nm in diameter, ␥precipitates are located uniformly (Fig. 2e,f) throughout the microstructure. It was not possible to confirm the presence of this phase by XRD analysis, due to the size and its small volume fraction of ␥. The microstructural stability of the hot rolled material was investigated by performing high temperature exposure for 168 h (7 days). LM analysis was used to investigated the grain size and TEM to analyze the precipitation phenomenon. The LM and TEM microstructures and corresponding selected area diffraction patterns are presented in Figs. 3–6. Additionally, in Fig. 6, SEM micrographs of the material after 900 ◦C exposure is presented. In the temperature range of 650–900 ◦C, the changes in the grain size are negligible. The growth of individual selected grains exceeding 100 ␮m can only be observed after exposure to 800 and 900 ◦C. After exposure to 900 ◦C, all grain boundaries are decorated by precipitates (Fig. 6a,c). Exposure to 650 ◦C caused the growth of the ␥precipitates, which existed in the material after rolling (Fig. 3b,c). In this state, the size of the ␥precipitates is between 50 and 200 nm. Fig. 2 – Micrographs of the investigated alloy in hot-rolled state. a) grains, LM, b) inverse pole figures, SEM-EBSD c) grains with twins, SEM-SE, d) twins, SEM-SE, e,f) small ␥precipitates, TEM. 554 j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 Fig. 3 – Micrographs after exposure to 650 ◦C for 168 h: a) grains, LM, b,c) ␥precipitates, TEM, and d) selective diffraction pattern of the area with ␥precipitates presented in c. Exposure to 700 ◦C caused significant changes. Continuous growth of existing after rolling ␥precipitates, additional continuous precipitation (CP) of ␥inside the grains (Fig. 4b,c) and discontinuous precipitation at the grain boundaries (Fig. 4b) occurred. Discontinuous precipitation (DP), also known as cellular precipitation, is a precipitation mode, which involves the cellular growth of alternating layers of the precipitated phase and matrix phase behind a moving high-angle grain boundary [19–21]. It can be seen in Fig. 4b that the DP structure occurred to a small extent along the grain boundaries, in the perpendicular direction. The grain boundaries were seen to migrate into a serrated shape due to the tendency of different parts of the same grain boundary to migrate in opposite directions. In our previous paper [22], where a short aging time (4 h) within the same temperature range was applied, we did not observe DP in the investigated alloy, even at a higher temperature, such as 900 ◦C. Based on literature reviews [21,23,24] and EDS analysis (Fig. 4e) it can be concluded that there were ␥precipitates interspersed with the ␥ matrix. ␥precipitates were additionally enriched in Ni, Co, Ti, and Al. Fe was located only in the ␥ matrix. According to the results presented in [21], DP observed in the investigated alloy may be a result of high cobalt contents. Based on Ni-Co-Al alloy investigations, the authors concluded that higher cobalt contents results in an increased susceptibility to DP during aging. Increasing the exposure temperature to 800 ◦C (Fig. 5) caused an increase in the number of areas where DP occurred, which differentiates HEA’s from simple Ni-Co-Al alloys [21,23], where DP usually appears during short time aging. Moreover, in the above-mentioned alloys, with an increase in temperature the susceptibility to CP was greater than DP and above 700 ◦C, DP did not occur. Authors of [25], suggest that DP will completely invade the microstructure only under conditions, in which the kinetics of CP are very slow. It seems that, as suggested in [26], there is still much work to be done to fully understand the mechanisms governing DP reactions in multicomponent systems. What is interesting, continuous precipitates of ␥inside the grains did not grow or the growth was negligible, which contradicts the results presented in [25]. Exposure to 900 ◦C (Fig. 6) promoted DP. All grain boundaries were decorated with DP with the length over 1 ␮m in the direction perpendicular to the boundaries. Additionally, DP areas were observed inside the grains (Fig. 6b,c). Probably, if the alloy was exposed for a longer duration, the entire volume of the material would consist of DP, which would be consistent with observation in work [21], where in Ni-Co-Al alloys after aging, DP would constituted throughout the entire volume of the materials. It is worth mentioning that ␥precipitates existing in the microstructure after hot rolling grew significantly (Fig. 6b). Some of them were larger than 500 nm, whereas continuous precipitates, which were created during exposure to 900 ◦C slightly increased in comparison to 700 ◦C (Fig. 6d). Finally, the microstructure of the investigated alloy consisted of ␥ matrix and three types of ␥precipitates, i.e. spherical ␥phase with bimodal distribution within the grains and elongated and DP ␥phase present at the grain boundaries. To confirm the above-mentioned changes in the microstructure caused by high-temperature exposure to the temperature range of 650–900 ◦C for 168 h the XRD analysis was performed (see Fig. 1). Only peaks related to ␥ (FCC solid solution) and the ␥phase (Ni3(Al,Ti)) were identified. The question is, how does the microstructure affect the mechanical properties of the alloy. The influence of the exposure temperature on hardness, yield strength (YS), and ultimate tensile strength (UTS) measured at 20 ◦C is presented j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 555 Fig. 4 – Micrographs after exposure to 700 ◦C for 168 h: a) grains, LM, b) discontinuous precipitation at the grain boundary, TEM, c) ␥precipitates, TEM, and d) selective diffraction pattern of the area with ␥precipitates presented in c, e) discontinuous precipitation at the grain boundary with EDS analysis, TEM. Table 2 – Mechanical properties of investigated alloy in as-rolled state and after high temperature exposure to the temperature range of 650–900 ◦C for 168 h. State YS [MPa] UTS [MPa] UTS/YS As-rolled 775 1045 1.348 650 ◦C/168 h 1050 1370 1.305 700 ◦C/168 h 990 1265 1.278 800 ◦C/168 h 705 1015 1.440 900 ◦C/168 h 490 770 1.571 in Fig. 7a and 7b. The data from the tensile tests are presented in Table 2. The highest hardness is exhibited by the sample exposed to 650 ◦C (455 HV10) and it exceeds the hardness in hot-rolled state by 130 HV. Moreover, CP of ␥’ observed after exposure to 700 ◦C did not cause an increase in hardness (430 HV10). This is related to the growth of ␥’ precipitates, which existed in the microstructure after hot rolling and the beginning of DP formation. It is consistent with the hardness decrease to 350 and 260 HV after the exposure of samples to 800 and 900 ◦C, respectively. The same relationship was observed in tensile test results. The material after hot rolling presents YS and UTS equal to 775 MPa and 1045 MPa, respectively. Exposure to 650 and 700 ◦C increased YS up to 1050 MPa and 990 MPa and UTS up to 1370 MPa and 1265 MPa, respectively. The application of higher exposure temperature decreased YS to 490 MPa and UTS to 770 MPa, with a significant increase of elongation for the material exposed at higher temperatures, i.e. ∼28% for sample exposed to 900 ◦C. 556 j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 Fig. 5 – Micrographs after exposure at 800 ◦C during 168 h: a) grains, LM, b) discontinuous precipitation at grain boundary, TEM c) ␥precipitates, TEM, and d) selective diffraction pattern of area with ␥precipitates presented in c. In the case of samples exposed to 650 or 700 ◦C, the high UTS is related to slip band accumulation in the final step of tensile tests, mainly on ␥precipitates, which were formed during cooling after hot rolling and have grown during high-temperature exposure. The ␥precipitates are used to strengthen nickel-based superalloys, where it is well established, so that a dislocation can penetrate the field of ␥precipitates by either bypassing (the so-called Orowan bypassing mechanism) or cutting the precipitates. The multimodal distribution of ␥allows superalloys to reach high yield strengths at temperatures even up to 750 ◦C, while simultaneously maintaining high creep and fatigue life properties [27,28]. In superalloys, which are based on the same elements, Kozar et al. [29], have proposed a model including multimodal ␥size effects by partially modifying the weak pair-coupling mechanism and introducing size distribution effects in the tertiary ␥. The highest strength is achieved when the secondary ␥particles are between 150 and 300 nm. After exposure to 650 ◦C, the size of the ␥precipitates in the investigated alloy is between 50 and 200 nm, which results in the highest strength. After exposure to 800 and 900 ◦C, the strength and hardness decreases as a result of continuing ␥particles growth (the ones, which precipitated during cooling after rolling), as dislocations no longer cut ␥particles, but bypass them. The ␥particles precipitated during exposure to 700 ◦C were too small (less than 50 nm) to play an important role. Based on the above-presented results we proposed the scheme of the microstructure evolution of the investigated alloy during hightemperature exposure and had an influence on mechanical properties of the alloy (Fig. 8). 4. Conclusions After hot rolling with air cooling the investigated alloy was not fully supersaturated. The application of air cooling resulted in the precipitation of ␥, which was uniformly distributed within the microstructure. Numerous aging twins and grain misorientations were observed in the microstructure. High-temperature exposure caused the growth of ␥particles. At 700 ◦C and above, additional CP and DP of ␥phase was observed. After applied heat treatment, the microstructure of the investigated alloy consisted of ␥ matrix and three types of ␥precipitates, i.e. spherical ␥phase with bimodal distribution within the grains and elongated, DP ␥phase present at the grain boundaries. Higher exposure temperatures tend to favor discontinuous precipitation. After exposure to 900 ◦C, all grain boundaries were decorated with DP with the length over 1 ␮m. Moreover, the size of ␥precipitated during air cooling after hot rolling increased significantly during high-temperature exposure. On the other hand, the difference in size of the ␥particles precipitated during exposure between 700 ◦C and 900 ◦C was negligible. Additionally, the grain size did not significantly change. These factors distinguish the investigated HEA’s from superalloys. The highest hardness and tensile properties were obtained after exposure to 650 ◦C, which was caused by ␥precipitates that precipitated during air cooling after hot rolling. Further increase in temperature decreased the mechanical properties. j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 557 Fig. 6 – Micrographs after exposure to 900 ◦C for 168 h: a) grains, LM, b) large ␥precipitates, SEM-SE c) discontinuous precipitation at the grain boundary, SEM-BSE d) continuous ␥precipitates inside grains, TEM, e) discontinuous precipitation at grain boundary, TEM and e) selective diffraction pattern of the area with ␥precipitates presented in d. Fig. 7 – a) hardness, and b) engineering stress-strain curves, obtained at room temperature in as-hot rolled state and after high-temperature exposure to the temperature range of 650–900 ◦C for 168 h. 558 j m a t e r r e s t e c h n o l . 2 0 2 0;9(1):551–559 Fig. 8 – Scheme of phase transformation pathway of Al5Ti5Co35Ni35Fe20 high-entropy alloy from as-rolled to annealed at 900 ◦C for 168 h, where DP-transformed regions gradually transform to near-equilibrium microstructure along with CP. Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgements This work was supported by the Polish National Science Centre [grant number 2017/25/N/ST8/02542]. r e f e r e n c e s [1] Callister WD Jr. Materials science and engineering: an introduction. USA: John Wiley & Sons; 2003. [2] Askeland DR, Wright WJ. The science and engineering of materials. 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