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Hot tensile properties of CoCrFeMnNi(NbC) compositionally complex alloys
Materials Science & Engineering A 772 (2020) 138771
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Hot tensile properties of CoCrFeMnNi(NbC) compositionally complex alloys Erfan Abbasi a, b, *, Kamran Dehghani a a b
Department of Materials and Metallurgical Engineering, Amirkabir University of Technology, Tehran, Iran Iran’s National Elites Foundation, Tehran, Iran
A R T I C L E I N F O
A B S T R A C T
Keywords: Compositionally complex alloys Nb–C addition Hot tensile testing Microstructure
An investigation was performed on tensile deformation behaviour and its relation with the microstructure of CoCrFeMnNi and CoCrFeMnNi(NbC) compositionally complex alloys at room temperature, 500, 600 and 700 � C. Microstructural characterisations and fractography were performed using optical microscopy, scanning electron microscopy and wavelength dispersive spectroscopy, X-ray diffraction, whereas hardness was obtained by Vickers hardness testing. The results showed that Nb–C addition can considerably increase the strength of CoCrFeMnNi alloy at room and high temperatures, while it can reduce the elongation up to about 50%. Mi croscopy observations demonstrated that the higher strength of CoCrFeMnNi(NbC) compared to CoCrFeMnNi was primarily related to the effect of carbide precipitates. Moreover, a larger grain size and carbide precipitates were responsible for the lower elongation of CoCrFeMnNi(NbC).
1. Introduction Generated heat due to friction or mechanical deformation as well as hot working environments has been a challenge in designing alloys for different engineering components (e.g. turbines, engines, exhaust sys tems etc) [1]. To meet such requirements, different alloys like superal loys, stainless/heat-resistant steels have been used for different applications [2,3]. In this way, High-Entropy-Alloys (HEA)/Composi tionally-Complex-Alloys (CCA) as a group of advanced materials were introduced in 2004 and have been considered as a promising alternative candidate for high temperature applications [1,4–6]. It should be mentioned that CCAs are a new generation of HEAs, but including sec ondary phases inside a high entropy matrix [7–9]. The secondary phases of CCAs can be a precipitate, ordered/disordered solid-solution phase and amorphous phase. These alloys have shown excellent cryogenic strength and ductility, outstanding thermal stability, good corrosion resistance and irradiation resistance [9–12]. However, CCAs are still new and they call for further investigations into their alloy design, processing parameters and mechanical properties. In HEAs, four factors are generally responsible for their outstanding mechanical/physical properties at working conditions, including high configurational entropy, severe lattice distortion, sluggish diffusion rate and cocktail effect (composite effect) [6,9]. The microstructure of HEAs consists of a high entropy (i.e. >1.5 R) solid solution matrix. The high entropy of solid solution matrix reduces the possibility of intermetallic
formation [6,13]. Besides, the presence of different alloying elements inside a lattice with different atomic radii causes severe lattice distortion which not only affect deformation mechanisms but also the performance of these materials [6,9]. The sluggish diffusion of HEA in addition to possible thermally stable precipitates in the microstructure of CCAs can act as a composite and enhance hot mechanical properties of HEAs [1, 14,15]. Nevertheless, very few works have been reported on the me chanical properties of CCAs at high temperatures [16–19]. CoCrFeMnNi (also known as Cantour HEA) has a stable single solid solution phase with outstanding mechanical properties at cryogenic and room temperatures [20,21]. However, further improvements in the strength of this alloy is still sought. Cantour HEA has been a good example to further develop the understanding about the alloy design of HEAs/CCAs. In this way, our previous studies and other reports in the literature on CoCrFeMnNi(NbC) have shown that Nb–C addition can significantly control the microstructure and enhance the strength of CoCrFeMnNi at room and cryogenic temperatures [22,23]. In spite of investigations into the high temperature mechanical properties of CoCrFeMnNi HEA [16,18,21], it is not still clear that how a thermo mechanically processed CoCrFeMnNi(NbC) behaves at high tempera tures (i.e. homologous temperatures of about 0.5–0.6 Tm). Therefore, in this investigation, the microstructure and mechanical properties of CoCrFeMnNi was studied against CoCrFeMnNi(NbC) at 500–700 � C to better understand the effect of Nb–C addition on hot tensile properties of such alloy systems.
* Corresponding author. Department of Mining and Metallurgical Engineering, Amirkabir University of Technology, Tehran, Iran. E-mail address: [email protected] (E. Abbasi). https://doi.org/10.1016/j.msea.2019.138771 Received 1 October 2019; Received in revised form 26 November 2019; Accepted 2 December 2019 Available online 3 December 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
E. Abbasi and K. Dehghani
Materials Science & Engineering A 772 (2020) 138771
with a gauge length of 6.9 mm and thickness of 1 mm. Tensile specimens were prepared by Electro Discharge Machining (EDM). The tensile specimen was designed according to JIS Z2201 and also tensile test was simulated by three-dimensional non-linear Finite Element (FE) method using Abaqus/Explicit commercial software (Dassault Syst� emes, Provi dence, Rhode Island) to ensure about the fitness of specimen geometry (Fig. 1 (b)). Isotropic properties and also two different testing temper atures (i.e. room temperature and 700 � C) were applied for the FE simulations. The testing conditions were considered isothermal in sim ulations and with a strain rate of 0.007 min 1. The repeatability of experimental results was checked using two separate samples from each alloy. Specimens were held 20 min immediately after reaching the target temperature and before the onset of tensile testing in order to stabilise the temperature throughout samples. The temperature of tensile spec imen was measured by three R-type thermocouples, attached to the lower, middle and upper section of specimen. The fractured specimens were rapidly air cooled from testing temperatures. The fractured surface and diffuse neck region of tensile samples were examined by optical and Scanning Electron Microscopy (SEM). The possible effect of holding time (before the onset of tensile testing) at target temperatures on the microstructure and hardness was checked by preparing separate samples. As-annealed samples with a size of 10x10 � 5 mm3 were heated to the target temperatures (i.e. 500, 600 and 700 � C) using a muffle furnace and were held for 20 min before cooling. This heat treatment is hereafter referred to as “heat treatment”. Microscopic characterisations and hardness measurements were performed on the normal direction plane of samples. Microstructure was studied by optical microscopy and SEM and qualitative Wavelengthdispersive X-ray spectroscopy (WDX) microanalysis. SEM analysis was
Table 1 Chemical composition of studied alloys (at%). Alloy
Co
Cr
Fe
Mn
Ni
Nb
C
CoCrFeMnNi(NbC) CoCrFeMnNi
19 19
20 20
20 20
21 21
19 19
0.06 –
0.8 –
2. Experimental procedure Two HEAs were melted from pure metals (purity > 99.8 wt%) by Vacuum Arc Re-melting (VAR) technique. Carbon powder was wrapped in pure Fe before melting to completely dissolve the carbon inside the melt. Also, an additional Mn (10 wt%) was added to each ingot to compensate for the evaporation due to its relatively high vapour pres sure. Re-melting and homogenisation at 1200 � C for 24 h (in argon) were carried out to enhance the homogeneity of chemical composition. Then, samples were subjected to cold rolling (i.e. 60% reduction) and subse quent annealing at 900 � C for 1 h, followed by air cooling. Table 1 gives the chemical composition of studied high-entropy alloys. The chemical composition was measured through XRF and the carbon content was determined by Leco-combustion technique. The standard deviation of XRF measurements for major elements was about �0.2 wt% and for minor element was �0.004 wt% and hence the values were rounded off to the nearest integer. The studied alloys were designated as “CoCr FeMnNi” and “CoCrFeMnNi(NbC)” in this paper. The tensile properties of annealed samples were evaluated using Shimadzu AG-25TC tensile tester. The tensile testing was carried out according to ASTM E21 at room temperature, 500, 600 and 700 � C (i.e. at homologous temperatures of ~0.5–0.6 Tm). Tensile testing was per formed using flat dog-bone shaped tensile specimens (rolling direction)
Fig. 1. (a) Tensile testing machine used in this study, (b) Illustration of tensile specimen (dimensions are in mm), (c) Abaqus simulation of tensile testing at 700 � C, showing effective plastic strain distribution along the designed specimen and within gauge length. 2
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Materials Science & Engineering A 772 (2020) 138771
Fig. 2. Optical micrographs, corresponding to the microstructure of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys before and after heat treatment at different temperatures.
done by a Philips XL 30 at 25 kV. The crystalline structure of samples after heat treatment was ana lysed by an X’Pert-Pro (PANalytical) X-ray diffractometer with Cu Kα radiation at 40 kV and 40 mA. The XRD analysis was carried out using the whole diffraction pattern through Rietveld refinement by the Topas Academic package software V5.0. Hardness measurements were performed by a universal Vickers hardness tester with 30 kg load and 10 s holding time before unloading.
The average hardness was calculated from five measurements per each sample. 3. Results 3.1. Microstructural evolution before the onset of tensile testing Fig. 2 presents the optical micrographs of studied alloys before and 3
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Materials Science & Engineering A 772 (2020) 138771
Fig. 3. SEM micrographs, corresponding to the microstructure of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys before and after heat treatment at different temperatures.
after heat treatment. The initial microstructure of CoCrFeMnNi alloy showed a much finer grain size (i.e. 22.1 � 1.4 μm) compared to CoCrFeMnNi(NbC) (i.e. over 100 μm), after identical thermomechanical processing. The results also indicated the absence of any significant grain refinement or growth/coarsening before the onset of tensile testing at all temperatures (i.e. 500–700 � C) in both alloys. A substruc ture was observed in optical micrographs of CoCrFeMnNi(NbC). Further SEM analyses revealed that these features are lath shaped structure surrounded by precipitates (Fig. 3). Besides, new distinctive features appeared in the optical micrographs of CoCrFeMnNi(NbC) after heat treatment at 600 and 700 � C (Fig. 2 (f) and (h)). SEM and XRD analyses of both alloys revealed the presence of Face Centred Cubic (FCC) single phase matrix before and after heat treatment (Figs. 3 and 5). XRD results also suggested insignificant difference
between the lattice parameter of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys after heat treatment (Table 2). SEM observations of microstructure showed the presence of pre cipitates in both alloys after heat treatment (Figs. 3 and 4). It was clear that precipitates were formed in both alloys before the onset of tensile testing. Other researchers have also reported the possibility of precipi tation in CoCrFeMnNi HEA at intermediate and high temperatures [24–27]. From the microstructural observations, it was also found that the density of precipitates in CoCrFeMnNi(NbC) was noticeably higher than CoCrFeMnNi alloy. In CoCrFeMnNi alloy, the precipitate size and density were rather similar in all samples and heat treatment tempera tures did not affect the results. Our previous theoretical study through JMatPro-V7.0 software (equilibrium conditions) suggested that only Nb-carbides and M23C6 4
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Materials Science & Engineering A 772 (2020) 138771
Fig. 4. Selected SEM micrographs, showing regions with a high density of precipitates in CoCrFeMnNi(NbC) alloy after the heat treatment at 600 � C.
3.2. Tensile testing FEM simulation evidenced an appropriate distribution of forces along the designed tensile specimens and also a concentrated plastic deformation within gauge length (Fig. 1 (c)). Experimental results also demonstrated that the fracture occurred inside the gauge length of tensile specimens (e.g. Fig. 1 (b)). This verified theoretical simulations and showed that the selected dimensions were suitable for the tensile testing. Fig. 6 compares engineering tensile curves of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys after testing at room temperature, 500, 600 and 700 � C. In general, the results showed a lower strength and lower total elongation due to temperature increases compared to room tem perature testing. These results agree with the findings of Otto et al. who also reported a similar behaviour for equiatomic CoCrFeMnNi [21]. The total elongation of CoCrFeMnNi(NbC) was considerably lower than CoCrFeMnNi (Fig. 6 (d)). Also it was found that the strength of CoCr FeMnNi(NbC) was generally higher than CoCrFeMnNi (Table 3). It was clear that the Nb–C addition could significantly reduce the total elon gation of CoCrFeMnNi at all testing temperatures. Nevertheless, the trend of strength variations at 600 and 700 � C could be different be tween the studied alloys. In both alloys, the yield strength and Ultimate Tensile Strength (UTS) were dramatically reduced at testing temperature of 500 � C (Table 3). By increasing the testing temperature to 600 and 700 � C the strength of CoCrFeMnNi alloy showed noticeable reductions. However, tensile re sults indicated that there was a mechanism that raised the strength of CoCrFeMnNi(NbC) at 600 and 700 � C, despite a higher testing temperature.
Fig. 5. Selected XRD spectrum, corresponding to CoCrFeMnNi (i.e. free) and CoCrFeMnNi(NbC) (i.e. NbC) alloys after heat treatment at different temperatures.
Table 2 Lattice parameter of heat treated samples, measured by XRD technique. CoCrFeMnNi CoCrFeMnNi(NbC)
Heat treatment temperature (� C)
Lattice parameter (nm)
500 600 700 500 600 700
0.3610 0.3611 0.3610 0.3611 0.3607 0.3609
3.3. Hardness
precipitates could form in the matrix of CoCrFeMnNi(NbC) [28]. Addi tionally, qualitative SEM-WDX analysis revealed the presence of Nb rich precipitates in the microstructure. Nevertheless, SEM-WDX results did not show any enrichment of other alloying elements in precipitates. This can be attributed to the relatively small size of these precipitates compared to the interaction volume of beam. Furthermore, SEM results showed that the observed features (as shown by arrows) in optical micrographs (Fig. 2) were related to the regions with a high density of precipitates in CoCrFeMnNi(NbC) (Fig. 4). Clearly, the density of precipitates in CoCrFeMnNi(NbC) alloy was significantly changed in terms of heat treatment temperature. By increasing the temperature up to 600 � C, the density of precipitates considerably and locally was raised so that regions with a relatively high density of precipitates were formed. However, as mentioned earlier, XRD results failed to show any secondary phases in the microstructure. This could be due to small volume fraction of these precipitates compared to the matrix.
Fig. 7 shows the hardness variations of studied alloys before and after heat treatments. The hardness of both studied alloys increased as a result of heat treatment, but with different trends. The hardness variations of CoCrFeMnNi was quite similar between heat treated samples at 500 and 600 � C, while a slight increase in hardness was observed after heat treating at 700 � C. In CoCrFeMnNi(NbC) alloy, the hardness slightly increased from 212 to 220 HV30 after heat treating at 500 � C. The hardness results also showed that the heat treatment at 600 � C could considerably increase the hardness of CoCrFeMnNi(NbC) to 288 HV30, whereas heat treatment at 700 � C was accompanied by a slight hardness reduction to 249 HV30. 3.4. Fractography of fractured specimens Fig. 8 highlights the fracture surface of tensile specimens after tensile testing at different temperatures. A comparison between CoCrFeMnNi and CoCrFeMnNi(NbC) alloys showed that the fracture behaviour of CoCrFeMnNi alloy could be changed at high temperatures due to Nb and 5
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Materials Science & Engineering A 772 (2020) 138771
Fig. 6. (a) and (b) Engineering stress-strain curves of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys, respectively, (c) Trend of yield and UTS variations, (d) Trend of elongation variations and elongation reduction (%) due to Nb–C addition (i.e. (%elCoCrFeMnNi – %elCoCrFeMnNi(NbC))/(%elCoCrFeMnNi)).
Table 3 Tensile properties of investigated alloys, tested at different temperatures. Alloy CoCrFeMnNi
CoCrFeMnNi (NbC)
Yield stress (MPa) UTS (MPa) Yield stress (MPa) UTS (MPa)
Troom
500 � C
600 � C
700 � C
433 � 9
242 13 407 11 246 15 377 11
150 10 307 14 368 15 480 12
141 � 11 209 � 16 259 � 9
726 � 12 500 � 8 696 � 10
� � � �
� � � �
289 � 13
C addition. A ductile dimple fracture appeared after testing at room temperature and 500 � C in CoCrFeMnNi alloy, while higher testing temperatures (i.e. 600 and 700 � C) enhanced an intergranular fracture. A dominant ductile dimple fracture was observed in CoCrFeMnNi(NbC) alloy at all testing temperatures. Additionally, the dimple size appeared to be different between studied alloys. To better understand the effect of testing temperature on the fracture mode of studied alloys, the microstructure of gauge length of fractured tensile specimens was also studied by optical microscopy. Fig. 9 shows selected microstructures of CoCrFeMnNi, corresponding to the diffuseneck-region/near-fracture-tip of fractured tensile specimens. Voids were seen along the grain boundaries and boundary junctions of tested CoCrFeMnNi specimens at 600 and 700 � C. However, such voids were not observed in other samples. Furthermore, no elongated grains were seen in the microstructure of deformed areas (Fig. 9). These observations are in line with the earlier discussed findings (Fig. 8 (e) and (g)),
Fig. 7. Hardness results of studied alloys, before and after the heat treatment at different temperatures.
6
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Fig. 8. SEM micrographs, showing the fracture surface of tensile specimens.
suggesting that grain boundary sliding occurred in CoCrFeMnNi alloy during tensile testing at 600 and 700 � C.
alloys. Our concurrent investigations into the annealing of these alloys have also shown that the annealing can reduce the strength of cold rolled samples, but with different rates [29]. It has been established that Nb–C addition could give higher strength to CoCrFeMnNi after cold rolling and annealing. It is thus clear that Nb–C addition increases the strength of CoCrFeMnNi HEA at room temperature from 433 � 9 to 500 � 8 MPa (Table 3). This is consistent with the results of Gao et al. who also re ported a similar behaviour for Nb–C added equiatomic CoCrFeMnNi HEA [23]. Moreover, our results ruled out the possible contribution of grain size to the yield stress, since CoCrFeMnNi with a smaller grain size (i.e. 22.1 � 1.4 μm) gave a lower strength compared to CoCrFeMnNi (NbC) with a larger grain size (i.e. over 100 μm). In this context, it has been shown by other researchers that the yield strength of CoCrFeMnNi HEA is slightly raised due to grain size refinement from ~150 to 20 μm [15,18,21]. The results showed the minimal effect of grain size on the
4. Discussion 4.1. Yield strength and hardness To interpret tensile testing results, the effect of microstructural evolution was taken into consideration, including microstructure and grain size of matrix as well as precipitation behaviour. The results showed that the yield stress and hardness of CoCrFeMnNi (NbC) were higher than CoCrFeMnNi at room temperature (Figs. 6 and 7). As the level of cold rolling reduction (i.e. 60%) before annealing was the same, the observed hardness difference stems from the prior cold rolling. In fact, cold rolling led to different work hardening in the studied 7
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Materials Science & Engineering A 772 (2020) 138771
Fig. 9. Optical micrographs, showing the formation of voids along boundaries/junctions in CoCrFeMnNi alloy during tensile testing, (a) At 600 � C, (b) At 700 � C.
strengthening of studied materials, as a larger grain size of CoCrFeMnNi (NbC) would have reduced the yield strength. Therefore, it is clear that the higher strength of CoCrFeMnNi(NbC) compared to CoCrFeMnNi was mainly attributed to the solid solution as well as precipitation strengthening of Nb–C. From the tensile analyses, it was also found that the yield strength was reduced in both alloys due to temperature increases, though the trend of variations was different between the studied alloys (Fig. 6). In CoCrFeMnNi, the yield strength was progressively reduced by increasing the temperature. However, the yield strength of CoCrFeMnNi(NbC) initially exhibited a reduction at testing temperature of 500 � C and it was again raised at 600 and 700 � C to a level far lower than room temperature yield strength. A comparison between two alloys demon strated that the yield tensile strength of both alloys was reduced and reached a similar level at 500 � C, i.e. ~240 MPa. These results also suggested that softening mechanisms were active in both alloys in spite of their different microstructures. In general, recovery and lower lattice friction stress due to an increase in temperature were likely the reasons for the strength reduction. This will be further clarified next. The observed increase in the yield strength of CoCrFeMnNi(NbC) at 600 and 700 � C suggested that active hardening mechanisms increased the strength of material against other competing softening mechanisms. An annealing induced hardening in HEAs has also been reported by other researchers [30–32]. They suggested that a hardening at temper atures over 700 � C can be mainly attributed to precipitation strength ening. In this case, the microstructural observations revealed the formation of regions with a high density of precipitates (Figs. 3 and 4). However, these features were not seen in CoCrFeMnNi alloy. According to the microscopy results, it can be inferred that the formation of such regions with a high density of precipitates was mainly responsible for the hardening of CoCrFeMnNi(NbC) at 600 and 700 � C. It should be mentioned that the hardness of CoCrFeMnNi alloy was also slightly raised after heat treatments at different applied tempera tures (Fig. 7). Other researchers also reported the possible hardening of equiatomic CoCrFeMnNi HEA at a similar range of temperature due to the formation of long-range order structure and precipitate strength ening [30,33–35]. Perhaps, the observed precipitates in CoCrFeMnNi HEA matrix could increase the hardness of material at room temperature (Fig. 3). However, tensile results indicated that these precipitates were not effective enough to increase the tensile yield strength at the high temperatures. Also, further investigations into other hardening mecha nisms such as long-range ordered structures in the studied alloys are suggested to better understand the observed phenomena.
(Fig. 6). Moreover, the results showed that the UTS of CoCrFeMnNi (NbC) remarkably increased at 600 and 700 � C, though this behaviour was not observed in CoCrFeMnNi. The formation of precipitates, in particular at 600 and 700 � C, was the most likely reason for the higher UTS of CoCrFeMnNi(NbC) (Figs. 3 and 4). In fact, the observed pre cipitates at grain interiors might deter dislocations motion and enhance work hardening. The results also suggested that the work hardening of CoCrFeMnNi (NbC) was significantly higher than CoCrFeMnNi at all temperatures (Fig. 6). Perhaps, the solid solution as well as precipitate strengthening due to Nb–C addition raised the work hardening of CoCrFeMnNi by the generation and multiplication of dislocations. 4.3. Total elongation The results evidenced that the total elongation of CoCrFeMnNi was almost twice as much as that of CoCrFeMnNi(NbC) at all temperatures. In addition, SEM observations of fractured specimens suggested that the fracture mechanism was different between studied alloys and at different testing temperatures. At room temperature and 500 � C, a dominant ductile fracture was observed in both alloys. Also, dimples appeared to be finer with a shallow depth in the fracture surface of CoCrFeMnNi(NbC) compared to CoCrFeMnNi. The observed difference was mainly attributed to the ef fect of grain size and precipitates that could accelerate fracture at room temperature and 500 � C. In fact, precipitates could enhance crack initiation, while a larger grain size promoted rapid crack propagation [14,36]. Moreover, Nb and C solute atoms could reduce the mobility of dislocations and consequently decrease the ductility. At 600 and 700 � C, the results also exhibited a higher elongation of CoCrFeMnNi compared to CoCrFeMnNi(NbC) (Fig. 6). Moreover, the fractured mode of CoCrFeMnNi(NbC) appeared to be a dominant ductile dimple fracture surface, while an intergranular fracture was observed in CoCrFeMnNi (Fig. 8). In this way, the microstructural studies of deformed regions (i.e. diffuse neck region) of fractured specimens also evidenced the formation of voids along grain boundaries and boundary junctions (Fig. 9). These results suggested that the intergranular fracture occurred due to grain boundary sliding. The possibility of grain boundary sliding has also been reported by other researchers for other materials and CoCrFeMnNi HEA at a similar range of temperature (i.e. at and above ~0.5Tm) [24,37–39]. They showed that the grain boundary sliding can be enhanced by refining the grain size. A comparison be tween the results of CoCrFeMnNi and CoCrFeMnNi(NbC) also suggested that a smaller grain size of CoCrFeMnNi (i.e. 22.1 � 1.4 μm) could enhance grain boundary sliding. Besides, the formation of voids in the microstructure of CoCrFeMnNi (Fig. 9) indicated that the grain bound ary sliding could accelerate the fracture, consequently reduced the elongation. Therefore, a slightly lower elongation of tested CoCrFeMnNi
4.2. Ultimate Tensile Strength The tensile testing results indicated that the UTS of CoCrFeMnNi (NbC) was higher than CoCrFeMnNi at room and high temperatures 8
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specimens at 600 and 700 � C compared to room temperature and 500 � C was mainly ascribed to this effect. However, a larger grain size and precipitates as well as solutes of Nb–C further reduced the elongation of CoCrFeMnNi(NbC) at 600–700 � C. As a whole, a lower elongation of CoCrFeMnNi(NbC) compared to CoCrFeMnNi was related to the effect of observed precipitates and Nb–C solute atoms. Precipitates and solute drag effect of Nb–C deterred dislocation motion and consequently supressed plastic deformation. Moreover, precipitates could enhance void nucleation and large grains accelerate rapid crack growth during fracture, promoting easier crack nucleation and growth [14,36].
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Hanold, E.P. George, Microstructural evolution after thermomechanical processing in an equiatomic, single-phase CoCrFeMnNi highentropy alloy with special focus on twin boundaries, Intermetallics 54 (2014) 39–48. [34] B. Schuh, F. Mendez-Martin, B. V€ olker, E.P. George, H. Clemens, R. Pippan, A. Hohenwarter, Mechanical properties, microstructure and thermal stability of a nanocrystalline CoCrFeMnNi high-entropy alloy after severe plastic deformation, Acta Mater. 96 (2015) 258–268. [35] J. Gu, S. Ni, Y. Liu, M. Song, Regulating the strength and ductility of a cold rolled FeCrCoMnNi high entropy alloy via annealing treatment, Mater. Sci. Eng., A 755 (2019) 289–294. [36] W. Li, P.K. Liaw, Y. Gao, Fracture resistance of high entropy alloys: a review, Intermetallics 99 (2018) 69–83. [37] H. Shahmir, J. Heb, Z. Lub, M. Kawasaki, T.G. Langdon, Evidence for superplasticity in a CoCrFeNiMn high-entropy alloy processed by high-pressure torsion, Mater. Sci. Eng., A 685 (2017) 342–348. [38] S.R. Reddy, S. Bapari, P.P. Bhattacharjee, A.H. Chokshi, Superplastic-like flow in a fine-grained equiatomic CoCrFeMnNi high-entropy alloy, Materials Research Letters 5 (6) (2017) 408–414. [39] T. Watanabe, Grain boundary design for desirable mechanical properties, J. Phys. Colloq. 49 (1988) 507–516.
5. Conclusions An investigation was carried out into hot tensile testing (i.e. 500–700 � C) of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys to better understand their tensile properties at room and high temperatures and their relation with microstructure. The following concluding remarks can be made: 1 Nb–C addition noticeably raised the tensile strengths of CoCrFeMnNi at all applied temperatures. This was mainly attributed to the for mation of carbide precipitates in the microstructure. 2 The studied alloys had two significantly different grain sizes, i.e. 22.1 � 1.4 and over 100 μm for CoCrFeMnNi and CoCrFeMnNi (NbC), respectively. The results showed that the grain size of both alloys was not changed during tensile testing, ruling out any soft ening due to grain-growth/coarsening and recrystallisation. 3 A ductile dimple fracture was dominant at room temperature and 500 � C in CoCrFeMnNi, while the fracture mode was altered to an intergranular fracture at 600 and 700 � C. The CoCrFeMnNi(NbC) also showed a ductile dimple fracture at all testing temperatures. The observed intergranular fracture was ascribed to grain boundary sliding. 4 Nb–C addition halved the total elongation of CoCrFeMnNi at all testing temperatures. The observed difference between the total elongation of studied alloys was primarily attributed to the effect of grain size and carbide precipitates. References [1] S. Praveen, H.S. Kim, “High-Entropy alloys: potential candidates for hightemperature applications – an overview, Adv. Eng. Mater. 20 (2018) 1700645. [2] H. Fecht, D. Furrer, Processing of nickel-base superalloys for turbine engine disc applications, Adv. Eng. Mater. 2 (2000) 777–787. [3] E.O. Ezugwu, Z.M. Wang, A.R. Machado, The machinability of nickel-based alloys: a review, J. Mater. Process. Technol. 86 (1999) 1–16. [4] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 375–377 (2004) 213–218. [5] E.P. George, D. Raabe, R.O. Ritchie, High-entropy alloys, Nature Reviews Materials 4 (2019) 515–534. [6] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1–93. [7] A.M. Manzoni, U. Glatzel, New multiphase compositionally complex alloys driven by the high entropy alloy approach, Mater. Char. 147 (2019) 512–532. [8] S. Gorsse, J.Ph Couzini� ec, D.B. Miracle, From high-entropy alloys to complex concentrated alloys, Compt. Rendus Phys. 19 (2018) 721–736. [9] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448–511. [10] E.J. Pickering, N.G. Jones, High-entropy alloys: a critical assessment of their founding principles and future prospects, Int. Mater. Rev. 61 (3) (2016) 183–202. [11] Y. Shi, B. Yang, P.K. Liaw, Corrosion-resistant high-entropy alloys: a review, Metals 7 (2) (2017) 43. [12] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, A fracture-resistant high-entropy alloy for cryogenic applications, Science 345 (6201) (2014) 1153–1158.
9
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Hot tensile properties of CoCrFeMnNi(NbC) compositionally complex alloys Erfan Abbasi a, b, *, Kamran Dehghani a a b
Department of Materials and Metallurgical Engineering, Amirkabir University of Technology, Tehran, Iran Iran’s National Elites Foundation, Tehran, Iran
A R T I C L E I N F O
A B S T R A C T
Keywords: Compositionally complex alloys Nb–C addition Hot tensile testing Microstructure
An investigation was performed on tensile deformation behaviour and its relation with the microstructure of CoCrFeMnNi and CoCrFeMnNi(NbC) compositionally complex alloys at room temperature, 500, 600 and 700 � C. Microstructural characterisations and fractography were performed using optical microscopy, scanning electron microscopy and wavelength dispersive spectroscopy, X-ray diffraction, whereas hardness was obtained by Vickers hardness testing. The results showed that Nb–C addition can considerably increase the strength of CoCrFeMnNi alloy at room and high temperatures, while it can reduce the elongation up to about 50%. Mi croscopy observations demonstrated that the higher strength of CoCrFeMnNi(NbC) compared to CoCrFeMnNi was primarily related to the effect of carbide precipitates. Moreover, a larger grain size and carbide precipitates were responsible for the lower elongation of CoCrFeMnNi(NbC).
1. Introduction Generated heat due to friction or mechanical deformation as well as hot working environments has been a challenge in designing alloys for different engineering components (e.g. turbines, engines, exhaust sys tems etc) [1]. To meet such requirements, different alloys like superal loys, stainless/heat-resistant steels have been used for different applications [2,3]. In this way, High-Entropy-Alloys (HEA)/Composi tionally-Complex-Alloys (CCA) as a group of advanced materials were introduced in 2004 and have been considered as a promising alternative candidate for high temperature applications [1,4–6]. It should be mentioned that CCAs are a new generation of HEAs, but including sec ondary phases inside a high entropy matrix [7–9]. The secondary phases of CCAs can be a precipitate, ordered/disordered solid-solution phase and amorphous phase. These alloys have shown excellent cryogenic strength and ductility, outstanding thermal stability, good corrosion resistance and irradiation resistance [9–12]. However, CCAs are still new and they call for further investigations into their alloy design, processing parameters and mechanical properties. In HEAs, four factors are generally responsible for their outstanding mechanical/physical properties at working conditions, including high configurational entropy, severe lattice distortion, sluggish diffusion rate and cocktail effect (composite effect) [6,9]. The microstructure of HEAs consists of a high entropy (i.e. >1.5 R) solid solution matrix. The high entropy of solid solution matrix reduces the possibility of intermetallic
formation [6,13]. Besides, the presence of different alloying elements inside a lattice with different atomic radii causes severe lattice distortion which not only affect deformation mechanisms but also the performance of these materials [6,9]. The sluggish diffusion of HEA in addition to possible thermally stable precipitates in the microstructure of CCAs can act as a composite and enhance hot mechanical properties of HEAs [1, 14,15]. Nevertheless, very few works have been reported on the me chanical properties of CCAs at high temperatures [16–19]. CoCrFeMnNi (also known as Cantour HEA) has a stable single solid solution phase with outstanding mechanical properties at cryogenic and room temperatures [20,21]. However, further improvements in the strength of this alloy is still sought. Cantour HEA has been a good example to further develop the understanding about the alloy design of HEAs/CCAs. In this way, our previous studies and other reports in the literature on CoCrFeMnNi(NbC) have shown that Nb–C addition can significantly control the microstructure and enhance the strength of CoCrFeMnNi at room and cryogenic temperatures [22,23]. In spite of investigations into the high temperature mechanical properties of CoCrFeMnNi HEA [16,18,21], it is not still clear that how a thermo mechanically processed CoCrFeMnNi(NbC) behaves at high tempera tures (i.e. homologous temperatures of about 0.5–0.6 Tm). Therefore, in this investigation, the microstructure and mechanical properties of CoCrFeMnNi was studied against CoCrFeMnNi(NbC) at 500–700 � C to better understand the effect of Nb–C addition on hot tensile properties of such alloy systems.
* Corresponding author. Department of Mining and Metallurgical Engineering, Amirkabir University of Technology, Tehran, Iran. E-mail address: [email protected] (E. Abbasi). https://doi.org/10.1016/j.msea.2019.138771 Received 1 October 2019; Received in revised form 26 November 2019; Accepted 2 December 2019 Available online 3 December 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Materials Science & Engineering A 772 (2020) 138771
with a gauge length of 6.9 mm and thickness of 1 mm. Tensile specimens were prepared by Electro Discharge Machining (EDM). The tensile specimen was designed according to JIS Z2201 and also tensile test was simulated by three-dimensional non-linear Finite Element (FE) method using Abaqus/Explicit commercial software (Dassault Syst� emes, Provi dence, Rhode Island) to ensure about the fitness of specimen geometry (Fig. 1 (b)). Isotropic properties and also two different testing temper atures (i.e. room temperature and 700 � C) were applied for the FE simulations. The testing conditions were considered isothermal in sim ulations and with a strain rate of 0.007 min 1. The repeatability of experimental results was checked using two separate samples from each alloy. Specimens were held 20 min immediately after reaching the target temperature and before the onset of tensile testing in order to stabilise the temperature throughout samples. The temperature of tensile spec imen was measured by three R-type thermocouples, attached to the lower, middle and upper section of specimen. The fractured specimens were rapidly air cooled from testing temperatures. The fractured surface and diffuse neck region of tensile samples were examined by optical and Scanning Electron Microscopy (SEM). The possible effect of holding time (before the onset of tensile testing) at target temperatures on the microstructure and hardness was checked by preparing separate samples. As-annealed samples with a size of 10x10 � 5 mm3 were heated to the target temperatures (i.e. 500, 600 and 700 � C) using a muffle furnace and were held for 20 min before cooling. This heat treatment is hereafter referred to as “heat treatment”. Microscopic characterisations and hardness measurements were performed on the normal direction plane of samples. Microstructure was studied by optical microscopy and SEM and qualitative Wavelengthdispersive X-ray spectroscopy (WDX) microanalysis. SEM analysis was
Table 1 Chemical composition of studied alloys (at%). Alloy
Co
Cr
Fe
Mn
Ni
Nb
C
CoCrFeMnNi(NbC) CoCrFeMnNi
19 19
20 20
20 20
21 21
19 19
0.06 –
0.8 –
2. Experimental procedure Two HEAs were melted from pure metals (purity > 99.8 wt%) by Vacuum Arc Re-melting (VAR) technique. Carbon powder was wrapped in pure Fe before melting to completely dissolve the carbon inside the melt. Also, an additional Mn (10 wt%) was added to each ingot to compensate for the evaporation due to its relatively high vapour pres sure. Re-melting and homogenisation at 1200 � C for 24 h (in argon) were carried out to enhance the homogeneity of chemical composition. Then, samples were subjected to cold rolling (i.e. 60% reduction) and subse quent annealing at 900 � C for 1 h, followed by air cooling. Table 1 gives the chemical composition of studied high-entropy alloys. The chemical composition was measured through XRF and the carbon content was determined by Leco-combustion technique. The standard deviation of XRF measurements for major elements was about �0.2 wt% and for minor element was �0.004 wt% and hence the values were rounded off to the nearest integer. The studied alloys were designated as “CoCr FeMnNi” and “CoCrFeMnNi(NbC)” in this paper. The tensile properties of annealed samples were evaluated using Shimadzu AG-25TC tensile tester. The tensile testing was carried out according to ASTM E21 at room temperature, 500, 600 and 700 � C (i.e. at homologous temperatures of ~0.5–0.6 Tm). Tensile testing was per formed using flat dog-bone shaped tensile specimens (rolling direction)
Fig. 1. (a) Tensile testing machine used in this study, (b) Illustration of tensile specimen (dimensions are in mm), (c) Abaqus simulation of tensile testing at 700 � C, showing effective plastic strain distribution along the designed specimen and within gauge length. 2
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Materials Science & Engineering A 772 (2020) 138771
Fig. 2. Optical micrographs, corresponding to the microstructure of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys before and after heat treatment at different temperatures.
done by a Philips XL 30 at 25 kV. The crystalline structure of samples after heat treatment was ana lysed by an X’Pert-Pro (PANalytical) X-ray diffractometer with Cu Kα radiation at 40 kV and 40 mA. The XRD analysis was carried out using the whole diffraction pattern through Rietveld refinement by the Topas Academic package software V5.0. Hardness measurements were performed by a universal Vickers hardness tester with 30 kg load and 10 s holding time before unloading.
The average hardness was calculated from five measurements per each sample. 3. Results 3.1. Microstructural evolution before the onset of tensile testing Fig. 2 presents the optical micrographs of studied alloys before and 3
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Materials Science & Engineering A 772 (2020) 138771
Fig. 3. SEM micrographs, corresponding to the microstructure of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys before and after heat treatment at different temperatures.
after heat treatment. The initial microstructure of CoCrFeMnNi alloy showed a much finer grain size (i.e. 22.1 � 1.4 μm) compared to CoCrFeMnNi(NbC) (i.e. over 100 μm), after identical thermomechanical processing. The results also indicated the absence of any significant grain refinement or growth/coarsening before the onset of tensile testing at all temperatures (i.e. 500–700 � C) in both alloys. A substruc ture was observed in optical micrographs of CoCrFeMnNi(NbC). Further SEM analyses revealed that these features are lath shaped structure surrounded by precipitates (Fig. 3). Besides, new distinctive features appeared in the optical micrographs of CoCrFeMnNi(NbC) after heat treatment at 600 and 700 � C (Fig. 2 (f) and (h)). SEM and XRD analyses of both alloys revealed the presence of Face Centred Cubic (FCC) single phase matrix before and after heat treatment (Figs. 3 and 5). XRD results also suggested insignificant difference
between the lattice parameter of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys after heat treatment (Table 2). SEM observations of microstructure showed the presence of pre cipitates in both alloys after heat treatment (Figs. 3 and 4). It was clear that precipitates were formed in both alloys before the onset of tensile testing. Other researchers have also reported the possibility of precipi tation in CoCrFeMnNi HEA at intermediate and high temperatures [24–27]. From the microstructural observations, it was also found that the density of precipitates in CoCrFeMnNi(NbC) was noticeably higher than CoCrFeMnNi alloy. In CoCrFeMnNi alloy, the precipitate size and density were rather similar in all samples and heat treatment tempera tures did not affect the results. Our previous theoretical study through JMatPro-V7.0 software (equilibrium conditions) suggested that only Nb-carbides and M23C6 4
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Fig. 4. Selected SEM micrographs, showing regions with a high density of precipitates in CoCrFeMnNi(NbC) alloy after the heat treatment at 600 � C.
3.2. Tensile testing FEM simulation evidenced an appropriate distribution of forces along the designed tensile specimens and also a concentrated plastic deformation within gauge length (Fig. 1 (c)). Experimental results also demonstrated that the fracture occurred inside the gauge length of tensile specimens (e.g. Fig. 1 (b)). This verified theoretical simulations and showed that the selected dimensions were suitable for the tensile testing. Fig. 6 compares engineering tensile curves of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys after testing at room temperature, 500, 600 and 700 � C. In general, the results showed a lower strength and lower total elongation due to temperature increases compared to room tem perature testing. These results agree with the findings of Otto et al. who also reported a similar behaviour for equiatomic CoCrFeMnNi [21]. The total elongation of CoCrFeMnNi(NbC) was considerably lower than CoCrFeMnNi (Fig. 6 (d)). Also it was found that the strength of CoCr FeMnNi(NbC) was generally higher than CoCrFeMnNi (Table 3). It was clear that the Nb–C addition could significantly reduce the total elon gation of CoCrFeMnNi at all testing temperatures. Nevertheless, the trend of strength variations at 600 and 700 � C could be different be tween the studied alloys. In both alloys, the yield strength and Ultimate Tensile Strength (UTS) were dramatically reduced at testing temperature of 500 � C (Table 3). By increasing the testing temperature to 600 and 700 � C the strength of CoCrFeMnNi alloy showed noticeable reductions. However, tensile re sults indicated that there was a mechanism that raised the strength of CoCrFeMnNi(NbC) at 600 and 700 � C, despite a higher testing temperature.
Fig. 5. Selected XRD spectrum, corresponding to CoCrFeMnNi (i.e. free) and CoCrFeMnNi(NbC) (i.e. NbC) alloys after heat treatment at different temperatures.
Table 2 Lattice parameter of heat treated samples, measured by XRD technique. CoCrFeMnNi CoCrFeMnNi(NbC)
Heat treatment temperature (� C)
Lattice parameter (nm)
500 600 700 500 600 700
0.3610 0.3611 0.3610 0.3611 0.3607 0.3609
3.3. Hardness
precipitates could form in the matrix of CoCrFeMnNi(NbC) [28]. Addi tionally, qualitative SEM-WDX analysis revealed the presence of Nb rich precipitates in the microstructure. Nevertheless, SEM-WDX results did not show any enrichment of other alloying elements in precipitates. This can be attributed to the relatively small size of these precipitates compared to the interaction volume of beam. Furthermore, SEM results showed that the observed features (as shown by arrows) in optical micrographs (Fig. 2) were related to the regions with a high density of precipitates in CoCrFeMnNi(NbC) (Fig. 4). Clearly, the density of precipitates in CoCrFeMnNi(NbC) alloy was significantly changed in terms of heat treatment temperature. By increasing the temperature up to 600 � C, the density of precipitates considerably and locally was raised so that regions with a relatively high density of precipitates were formed. However, as mentioned earlier, XRD results failed to show any secondary phases in the microstructure. This could be due to small volume fraction of these precipitates compared to the matrix.
Fig. 7 shows the hardness variations of studied alloys before and after heat treatments. The hardness of both studied alloys increased as a result of heat treatment, but with different trends. The hardness variations of CoCrFeMnNi was quite similar between heat treated samples at 500 and 600 � C, while a slight increase in hardness was observed after heat treating at 700 � C. In CoCrFeMnNi(NbC) alloy, the hardness slightly increased from 212 to 220 HV30 after heat treating at 500 � C. The hardness results also showed that the heat treatment at 600 � C could considerably increase the hardness of CoCrFeMnNi(NbC) to 288 HV30, whereas heat treatment at 700 � C was accompanied by a slight hardness reduction to 249 HV30. 3.4. Fractography of fractured specimens Fig. 8 highlights the fracture surface of tensile specimens after tensile testing at different temperatures. A comparison between CoCrFeMnNi and CoCrFeMnNi(NbC) alloys showed that the fracture behaviour of CoCrFeMnNi alloy could be changed at high temperatures due to Nb and 5
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Fig. 6. (a) and (b) Engineering stress-strain curves of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys, respectively, (c) Trend of yield and UTS variations, (d) Trend of elongation variations and elongation reduction (%) due to Nb–C addition (i.e. (%elCoCrFeMnNi – %elCoCrFeMnNi(NbC))/(%elCoCrFeMnNi)).
Table 3 Tensile properties of investigated alloys, tested at different temperatures. Alloy CoCrFeMnNi
CoCrFeMnNi (NbC)
Yield stress (MPa) UTS (MPa) Yield stress (MPa) UTS (MPa)
Troom
500 � C
600 � C
700 � C
433 � 9
242 13 407 11 246 15 377 11
150 10 307 14 368 15 480 12
141 � 11 209 � 16 259 � 9
726 � 12 500 � 8 696 � 10
� � � �
� � � �
289 � 13
C addition. A ductile dimple fracture appeared after testing at room temperature and 500 � C in CoCrFeMnNi alloy, while higher testing temperatures (i.e. 600 and 700 � C) enhanced an intergranular fracture. A dominant ductile dimple fracture was observed in CoCrFeMnNi(NbC) alloy at all testing temperatures. Additionally, the dimple size appeared to be different between studied alloys. To better understand the effect of testing temperature on the fracture mode of studied alloys, the microstructure of gauge length of fractured tensile specimens was also studied by optical microscopy. Fig. 9 shows selected microstructures of CoCrFeMnNi, corresponding to the diffuseneck-region/near-fracture-tip of fractured tensile specimens. Voids were seen along the grain boundaries and boundary junctions of tested CoCrFeMnNi specimens at 600 and 700 � C. However, such voids were not observed in other samples. Furthermore, no elongated grains were seen in the microstructure of deformed areas (Fig. 9). These observations are in line with the earlier discussed findings (Fig. 8 (e) and (g)),
Fig. 7. Hardness results of studied alloys, before and after the heat treatment at different temperatures.
6
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Materials Science & Engineering A 772 (2020) 138771
Fig. 8. SEM micrographs, showing the fracture surface of tensile specimens.
suggesting that grain boundary sliding occurred in CoCrFeMnNi alloy during tensile testing at 600 and 700 � C.
alloys. Our concurrent investigations into the annealing of these alloys have also shown that the annealing can reduce the strength of cold rolled samples, but with different rates [29]. It has been established that Nb–C addition could give higher strength to CoCrFeMnNi after cold rolling and annealing. It is thus clear that Nb–C addition increases the strength of CoCrFeMnNi HEA at room temperature from 433 � 9 to 500 � 8 MPa (Table 3). This is consistent with the results of Gao et al. who also re ported a similar behaviour for Nb–C added equiatomic CoCrFeMnNi HEA [23]. Moreover, our results ruled out the possible contribution of grain size to the yield stress, since CoCrFeMnNi with a smaller grain size (i.e. 22.1 � 1.4 μm) gave a lower strength compared to CoCrFeMnNi (NbC) with a larger grain size (i.e. over 100 μm). In this context, it has been shown by other researchers that the yield strength of CoCrFeMnNi HEA is slightly raised due to grain size refinement from ~150 to 20 μm [15,18,21]. The results showed the minimal effect of grain size on the
4. Discussion 4.1. Yield strength and hardness To interpret tensile testing results, the effect of microstructural evolution was taken into consideration, including microstructure and grain size of matrix as well as precipitation behaviour. The results showed that the yield stress and hardness of CoCrFeMnNi (NbC) were higher than CoCrFeMnNi at room temperature (Figs. 6 and 7). As the level of cold rolling reduction (i.e. 60%) before annealing was the same, the observed hardness difference stems from the prior cold rolling. In fact, cold rolling led to different work hardening in the studied 7
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Materials Science & Engineering A 772 (2020) 138771
Fig. 9. Optical micrographs, showing the formation of voids along boundaries/junctions in CoCrFeMnNi alloy during tensile testing, (a) At 600 � C, (b) At 700 � C.
strengthening of studied materials, as a larger grain size of CoCrFeMnNi (NbC) would have reduced the yield strength. Therefore, it is clear that the higher strength of CoCrFeMnNi(NbC) compared to CoCrFeMnNi was mainly attributed to the solid solution as well as precipitation strengthening of Nb–C. From the tensile analyses, it was also found that the yield strength was reduced in both alloys due to temperature increases, though the trend of variations was different between the studied alloys (Fig. 6). In CoCrFeMnNi, the yield strength was progressively reduced by increasing the temperature. However, the yield strength of CoCrFeMnNi(NbC) initially exhibited a reduction at testing temperature of 500 � C and it was again raised at 600 and 700 � C to a level far lower than room temperature yield strength. A comparison between two alloys demon strated that the yield tensile strength of both alloys was reduced and reached a similar level at 500 � C, i.e. ~240 MPa. These results also suggested that softening mechanisms were active in both alloys in spite of their different microstructures. In general, recovery and lower lattice friction stress due to an increase in temperature were likely the reasons for the strength reduction. This will be further clarified next. The observed increase in the yield strength of CoCrFeMnNi(NbC) at 600 and 700 � C suggested that active hardening mechanisms increased the strength of material against other competing softening mechanisms. An annealing induced hardening in HEAs has also been reported by other researchers [30–32]. They suggested that a hardening at temper atures over 700 � C can be mainly attributed to precipitation strength ening. In this case, the microstructural observations revealed the formation of regions with a high density of precipitates (Figs. 3 and 4). However, these features were not seen in CoCrFeMnNi alloy. According to the microscopy results, it can be inferred that the formation of such regions with a high density of precipitates was mainly responsible for the hardening of CoCrFeMnNi(NbC) at 600 and 700 � C. It should be mentioned that the hardness of CoCrFeMnNi alloy was also slightly raised after heat treatments at different applied tempera tures (Fig. 7). Other researchers also reported the possible hardening of equiatomic CoCrFeMnNi HEA at a similar range of temperature due to the formation of long-range order structure and precipitate strength ening [30,33–35]. Perhaps, the observed precipitates in CoCrFeMnNi HEA matrix could increase the hardness of material at room temperature (Fig. 3). However, tensile results indicated that these precipitates were not effective enough to increase the tensile yield strength at the high temperatures. Also, further investigations into other hardening mecha nisms such as long-range ordered structures in the studied alloys are suggested to better understand the observed phenomena.
(Fig. 6). Moreover, the results showed that the UTS of CoCrFeMnNi (NbC) remarkably increased at 600 and 700 � C, though this behaviour was not observed in CoCrFeMnNi. The formation of precipitates, in particular at 600 and 700 � C, was the most likely reason for the higher UTS of CoCrFeMnNi(NbC) (Figs. 3 and 4). In fact, the observed pre cipitates at grain interiors might deter dislocations motion and enhance work hardening. The results also suggested that the work hardening of CoCrFeMnNi (NbC) was significantly higher than CoCrFeMnNi at all temperatures (Fig. 6). Perhaps, the solid solution as well as precipitate strengthening due to Nb–C addition raised the work hardening of CoCrFeMnNi by the generation and multiplication of dislocations. 4.3. Total elongation The results evidenced that the total elongation of CoCrFeMnNi was almost twice as much as that of CoCrFeMnNi(NbC) at all temperatures. In addition, SEM observations of fractured specimens suggested that the fracture mechanism was different between studied alloys and at different testing temperatures. At room temperature and 500 � C, a dominant ductile fracture was observed in both alloys. Also, dimples appeared to be finer with a shallow depth in the fracture surface of CoCrFeMnNi(NbC) compared to CoCrFeMnNi. The observed difference was mainly attributed to the ef fect of grain size and precipitates that could accelerate fracture at room temperature and 500 � C. In fact, precipitates could enhance crack initiation, while a larger grain size promoted rapid crack propagation [14,36]. Moreover, Nb and C solute atoms could reduce the mobility of dislocations and consequently decrease the ductility. At 600 and 700 � C, the results also exhibited a higher elongation of CoCrFeMnNi compared to CoCrFeMnNi(NbC) (Fig. 6). Moreover, the fractured mode of CoCrFeMnNi(NbC) appeared to be a dominant ductile dimple fracture surface, while an intergranular fracture was observed in CoCrFeMnNi (Fig. 8). In this way, the microstructural studies of deformed regions (i.e. diffuse neck region) of fractured specimens also evidenced the formation of voids along grain boundaries and boundary junctions (Fig. 9). These results suggested that the intergranular fracture occurred due to grain boundary sliding. The possibility of grain boundary sliding has also been reported by other researchers for other materials and CoCrFeMnNi HEA at a similar range of temperature (i.e. at and above ~0.5Tm) [24,37–39]. They showed that the grain boundary sliding can be enhanced by refining the grain size. A comparison be tween the results of CoCrFeMnNi and CoCrFeMnNi(NbC) also suggested that a smaller grain size of CoCrFeMnNi (i.e. 22.1 � 1.4 μm) could enhance grain boundary sliding. Besides, the formation of voids in the microstructure of CoCrFeMnNi (Fig. 9) indicated that the grain bound ary sliding could accelerate the fracture, consequently reduced the elongation. Therefore, a slightly lower elongation of tested CoCrFeMnNi
4.2. Ultimate Tensile Strength The tensile testing results indicated that the UTS of CoCrFeMnNi (NbC) was higher than CoCrFeMnNi at room and high temperatures 8
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specimens at 600 and 700 � C compared to room temperature and 500 � C was mainly ascribed to this effect. However, a larger grain size and precipitates as well as solutes of Nb–C further reduced the elongation of CoCrFeMnNi(NbC) at 600–700 � C. As a whole, a lower elongation of CoCrFeMnNi(NbC) compared to CoCrFeMnNi was related to the effect of observed precipitates and Nb–C solute atoms. Precipitates and solute drag effect of Nb–C deterred dislocation motion and consequently supressed plastic deformation. Moreover, precipitates could enhance void nucleation and large grains accelerate rapid crack growth during fracture, promoting easier crack nucleation and growth [14,36].
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5. Conclusions An investigation was carried out into hot tensile testing (i.e. 500–700 � C) of CoCrFeMnNi and CoCrFeMnNi(NbC) alloys to better understand their tensile properties at room and high temperatures and their relation with microstructure. The following concluding remarks can be made: 1 Nb–C addition noticeably raised the tensile strengths of CoCrFeMnNi at all applied temperatures. This was mainly attributed to the for mation of carbide precipitates in the microstructure. 2 The studied alloys had two significantly different grain sizes, i.e. 22.1 � 1.4 and over 100 μm for CoCrFeMnNi and CoCrFeMnNi (NbC), respectively. The results showed that the grain size of both alloys was not changed during tensile testing, ruling out any soft ening due to grain-growth/coarsening and recrystallisation. 3 A ductile dimple fracture was dominant at room temperature and 500 � C in CoCrFeMnNi, while the fracture mode was altered to an intergranular fracture at 600 and 700 � C. The CoCrFeMnNi(NbC) also showed a ductile dimple fracture at all testing temperatures. The observed intergranular fracture was ascribed to grain boundary sliding. 4 Nb–C addition halved the total elongation of CoCrFeMnNi at all testing temperatures. The observed difference between the total elongation of studied alloys was primarily attributed to the effect of grain size and carbide precipitates. References [1] S. Praveen, H.S. Kim, “High-Entropy alloys: potential candidates for hightemperature applications – an overview, Adv. Eng. Mater. 20 (2018) 1700645. [2] H. Fecht, D. Furrer, Processing of nickel-base superalloys for turbine engine disc applications, Adv. Eng. Mater. 2 (2000) 777–787. [3] E.O. Ezugwu, Z.M. Wang, A.R. Machado, The machinability of nickel-based alloys: a review, J. Mater. Process. Technol. 86 (1999) 1–16. [4] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 375–377 (2004) 213–218. [5] E.P. George, D. Raabe, R.O. Ritchie, High-entropy alloys, Nature Reviews Materials 4 (2019) 515–534. [6] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1–93. [7] A.M. Manzoni, U. Glatzel, New multiphase compositionally complex alloys driven by the high entropy alloy approach, Mater. Char. 147 (2019) 512–532. [8] S. Gorsse, J.Ph Couzini� ec, D.B. Miracle, From high-entropy alloys to complex concentrated alloys, Compt. Rendus Phys. 19 (2018) 721–736. [9] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448–511. [10] E.J. Pickering, N.G. Jones, High-entropy alloys: a critical assessment of their founding principles and future prospects, Int. Mater. Rev. 61 (3) (2016) 183–202. [11] Y. Shi, B. Yang, P.K. Liaw, Corrosion-resistant high-entropy alloys: a review, Metals 7 (2) (2017) 43. [12] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, A fracture-resistant high-entropy alloy for cryogenic applications, Science 345 (6201) (2014) 1153–1158.
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