Journal of Materials Science & Technology  2020 , 38 (0): 19-27 https://doi.org/10.1016/j.jmst.2019.08.019

Research Article

Microstructure evolution, Cu segregation and tensile properties of CoCrFeNiCu high entropy alloy during directional solidification

Huiting Zhenga, Ruirun Chenab*, Gang Qina, Xinzhong Lia, Yanqing Suab, Hongsheng Dinga, Jingjie Guoa, Hengzhi Fua

aNational Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, China
bState Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China

Corresponding authors:   ∗Corresponding author at: National Key Laboratory for Precision Hot Processingof Metals, Harbin Institute of Technology, Harbin 150001, China. E-mail address: ruirunchen@hit.edu.cn (R. Chen).

Received: 2019-03-9

Revised:  2019-05-31

Accepted:  2019-08-7

Online:  2020-02-01

Copyright:  2020 Editorial board of Journal of Materials Science & Technology Copyright reserved, Editorial board of Journal of Materials Science & Technology

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Abstract

CoCrFeNiCu (equiatomic ratio) samples (ø 8 mm) were directionally solidified at different velocities (10, 30, 60 and 100 μm/s) to investigate the relationship between solidification velocity and microstructure formation, Cu micro-segregation as well as tensile properties. The results indicate that the morphology of the solid-liquid (S-L) interface evolves from convex to planar and then to concave with the increase of solidification velocity. Meanwhile, the primary and the secondary dendritic arm spacings decrease from 100 μm to 10 μm and from 20 μm to 5 μm, respectively. They are mainly influenced by the axial heat transfer and grain competition growth. During directional solidification, element Cu is repelled from the FCC phase and accumulates in the liquid owe to its positive mixing enthalpy with other elements. Tensile testing results show that the ultimate tensile strength (UTS) gradually increases from 400 MPa to 450 MPa, and the strain of the specimen prepared at the velocity of 60 μm/s is higher than those of others. The fracture mode of all specimens is the mixed fracture containing both ductile fracture and brittle fracture, in which ductile fracture plays a fundamental role. In addition, the brittle fracture is induced by Cu segregation. The improvement of UTS is resulted from columnar grain boundary strengthening.

Keywords: High entropy alloy ; Directional solidification ; Segregation ; Tensile properties

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Huiting Zheng, Ruirun Chen, Gang Qin, Xinzhong Li, Yanqing Su, Hongsheng Ding, Jingjie Guo, Hengzhi Fu. Microstructure evolution, Cu segregation and tensile properties of CoCrFeNiCu high entropy alloy during directional solidification[J]. Journal of Materials Science & Technology, 2020, 38(0): 19-27 https://doi.org/10.1016/j.jmst.2019.08.019

1. Introduction

High entropy alloys (HEAs) are new advanced alloys which contain more than five different elements with equal or near-equal molar contents [[1], [2], [3], [4], [5], [6], [7]]. Comparing with traditional alloys, HEAs show the relatively superior properties in some respects, such as high strength and corrosion resistance as well as the outstanding ductility for their special elemental content. Therefore, HEAs are identified as appropriate candidates for structural materials [[8], [9], [10]]. The microstructures of HEAs usually consist of FCC or BCC solid solution phase rather than the typical ordered phase due to the high entropy effect [[11], [12], [13]]. Nevertheless, the relationship between the microstructure evolution and mechanical property is still not clear, and it has become research hotspot in the development of HEAs.

Microstructure evolutions of as-cast HEAs are often investigated by adding different elements or controlling their contents [[14], [15], [16]]. However, the important parameters during solidification, such as solidification velocity (v) and temperature gradient (G) can not be controlled easily. Therefore, the microstructure formation and evolution are uncontrollable [[17], [18], [19], [20]]. A new preparing method is urgently needed to be developed to control the solidification parameters. Directional solidification technique is an effective method to control the solidification velocity, temperature gradient and the morphology of S-L interface during solidification [[21], [22], [23]]. Furthermore, it can eliminate (or restrict) the horizontal grain boundary of the alloys, which is considered as an important method to improve mechanical properties. Therefore, directional solidification is a promising method for the microstructure evolution of HEAs.

CoCrFeNiCu HEA has been one of the most extensively studied HEAs due to its excellent comprehensive performance. Nevertheless, the positive mixing enthalpy between Cu and Co, Cr, Fe and the rather small negative mixing enthalpy between Cu and Ni lead to serious Cu segregation and casting defects [[24], [25], [26], [27]]. Therefore, many researchers tried to reduce or eliminate Cu segregation by decreasing the atomic ratio of Cu, thus improving the microstructure and mechanical properties [[28], [29], [30], [31]]. However, the control of Cu segregation in equiatomic CoCrFeNiCu HEA during solidification is seldom reported. More importantly, the relationship between Cu segregation and solidification parameters as well as between microstructure evolution and mechanical properties of CoCrFeNiCu HEA are still not clear.

In this study, the equimolar CoCrFeNiCu HEA was prepared by directional solidification at different solidification velocities. Microstructure evolution, Cu segregation and tensile properties were systematically investigated to reveal the relationship between solidification parameters and microstructure evolution, Cu segregation as well as tensile properties.

2. Experimental details

Master ingots of equimolar CoCrFeNiCu HEA were prepared by the arc melting furnace. Each ingot was re-melted 5 times to ensure compositional homogeneity. Then the vacuum suction casting process was applied to drop the liquid alloy into a cylindrical rod mold with the diameter of 8 mm.

CoCrFeNiCu samples were directionally solidified in the vacuum electromagnetic induction furnace. This furnace consists of vacuum system, heating system, pulling system and cooling system. The cooling system contains the recirculation cooling water and Ga-In-Sn liquid alloy which is used to cool the new samples. CoCrFeNiCu sample was placed in the alumina tube, and they were surrounded by a graphite sheath. The graphite sheath was heated by induction coil, and the sample was heated by the radiant heat transfer from the graphite. To ensure temperature homogenization, the liquid alloy was held for 60 min after the rod was melted. Fresh samples were directionally solidified at the velocities of 10, 30, 60 and 100 μm/s, and the length of them is about 75 mm, as shown in Fig. 1. In order to preserve the morphology of S-L interface, the samples were quenched in the Ga-In-Sn liquid alloy. In this research, the pulling velocity of the sample can be approximately equal to the solidification velocity.

Fig. 1.   Directionally solidified CoCrFeNiCu HEA sample.

The directionally solidified samples were cut along longitudinal direction, and the microstructure was observed by optical microscopy (OM) and scanning electron microscopy (SEM). Energy dispersive spectroscopy (EDS) was carried out to investigate the element distribution of different phases. In order to further investigate the Cu segregation of CoCrFeNiCu HEA, the electron probe micro-analyzer (EPMA) area scan was also utilized to test the elements distribution. Tensile properties were carried out on a Zwick 150 type universal tensile testing machine, and the test rate was 1 mm/min. The tensile rod specimen was machined from the directionally solidified sample, and the gauge length and diameter of tensile specimen are 20 mm and 4 mm, respectively.

3. Results and discussion

3.1. Effect of solidification velocity on the morphology of solid-liquid interface

Fig. 2 shows the S-L interface of CoCrFeNiCu HEA at the solidification velocity of 10 μm/s. Mushy zone, dendritic tip and the inter-dendritic region are shown in Fig. 2(b)-(d), respectively. Equiaxed grains with size smaller than 10 μm were formed in the front of the dendritic tip, as shown in Fig. 2(b) and (c). The formation of these equiaxed grains is resulted from the quenching of sample into the liquid Ga-In-Sn alloy. The orientations of some grains are perpendicular to the heat flow, and some are parallel to the heat flow. This result proves that the growing direction of the primary dendrites is along the direction of heat flow and that of the secondary dendrites is perpendicular to the direction of the heat flow when the solidification velocity is slower. The equiaxed grains with smaller size about 5 μm are distributed in the inter-dendritic regions, as shown in Fig. 2(d).

Fig. 2.   S-L interface of CoCrFeNiCu HEA directionally solidified at 10 μm/s: (a) S-L interface; (b) mushy zone; (c) dendrite and dendritic tip; (d) inter-dendritic region.

It is well-known that microstructure evolution is significantly influenced by solidification parameters such as v and G. CoCrFeNiCu HEA samples were directionally solidified at different solidification velocities (v = 10-100 μm/s) without changing the temperature gradient, the morphologies of S-L interface of them are shown in Fig. 3. The primary and secondary dendritic arm spacing of the solidified samples are calculated and shown in Fig. 4. When the solidification velocity is 10 μm/s, the morphology of the S-L interface is convex (as shown in Fig. 3(a)). The primary dendritic arm spacing is 100 μm (as shown in Fig. 4(a)) and the secondary dendritic spacing arm is 20 μm (as shown in Fig. 4(b)). When the solidification velocity increases to 30 μm/s, the morphology of the S-L interface becomes planar. The primary dendritic arm spacing is 50 μm and the secondary dendritic arm spacing is 10 μm. When the solidification velocity further increases to 60 μm/s, the morphology of the S-L interface remains planar. Nonetheless, the primary dendritic arm spacing decreases to 25 μm and the secondary dendritic arm spacing decreases to 6 μm. When the solidification velocity increases to 100 μm/s, the morphology of the S-L interface evolves to concave shape. The primary dendritic and secondary dendritic arm spacing are 10 μm and 5 μm, respectively.

Fig. 3.   Solidification front of CoCrFeNiCu HEA at different solidification velocities: (a) 10 μm/s; (b) 30 μm/s; (c) 60 μm/s; (d) 100 μm/s.

Fig. 4.   Dendritic arm spacing of the samples prepared at different velocities: (a) the primary dendritic arm spacing; (b) the secondary dendritic arm spacing.

With the increase of solidification velocity, the morphology of the S-L interface evolves from convex to planar and then to concave, this result could be explained by the balance of heat input and output during solidification. As shown in Fig. 5, when the solidification velocity is slow (10 μm/s), the heat input is much larger than the heat output. The heat input in both sides of the S-L interface is apparently larger than that in the center of the S-L interface during directional solidification. In other words, grains in both sides of the S-L interface are easily to be re-melted because the heat input is larger than the heat output in this region, which results in the convex morphology of the S-L interface. When the solidification velocity increases to 30 and 60 μm/s, the heat input and output could reach a relative balance, which means that the temperature field in both sides and the center of the S-L interface is uniform. The uniform temperature field results in the growth of columnar grains and therefore planar S-L interface. When the solidification velocity further increases to 100 μm/s, the heat output is larger than the heat input in unit time, this leads to the temperatures in both sides are lower than that in the center. The outer layer of the sample (both sides of S-L interface) will be cooled faster when the sample is pulled into the Ga-In-Sn liquid alloy at the higher velocity, the growth rates of the grains on both sides of S-L interface are higher than that in the center. Therefore, the morphology of the S-L interface evolves from planar to concave.

Fig. 5.   Evolution of S-L interface during directional solidification.

3.2. The liquid phase separation of CoCrFeNiCu HEA during directional solidification

Previous study indicates that the liquidus and solidus temperatures of CoCrFeNiCu HEA are 1655 K and 1380 K, respectively, and XRD result shows that CoCrFeNiCu HEA consists of CoCrFeNiCu FCC solid solution phase and Cu-enriched solid solution phase [32]. To further study the phase formation of the CoCrFeNiCu HEA during directional solidification, chemical compositions of different positions in the sample directionally solidified at 10 μm/s were examined by EDS, the testing positions are shown in Fig. 6 and the EDS results are shown in Table 1.

Fig. 6.   Testing positions by EDS in CoCrFeNiCu HEA directionally solidified at velocity 10 μm/s: (a) solidification front; (b) mushy zone. A-inter-dendritic region; B-dendrite; C-grains in the mushy zone; D-grain boundary in the mushy zone.

Table 1   Chemical composition (at.%) of directionally solidified CoCrFeNiCu HEA.

PositionCuCrFeCoNi
A76.684.644.043.6810.95
B16.4121.2223.2619.5419.56
C9.6324.2622.7623.5219.84
D70.96.916.485.6210.09

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The EDS results indicate that the element segregation in the inter-dendritic region is serious. The content of Cu is 76.68 at.%, the content of Co, Cr and Fe are all less than 5 at.%, and the content of Ni is about 10 at.%. It means that the contents of the elements in the inter-dendritic region are lower than their average content except element Cu. The contents of Co, Cr, Ni and Fe in the dendritic region (marked as B in Fig. 6(a)) are about 20 at.%, they are close to their nominal composition, but the content of Cu is less than 20 at.%. It can be seen that there is a big difference in the element distribution of the inter-dendritic and dendritic regions. The EDS results show there still exists serious Cu element segregation in the mushy zone. Element Cu is enriched in the grain boundaries, while the content of other elements in the grains are close to their nominal compositions.

During directional solidification, microstructure evolution and element segregation are significantly affected by the diffusion and convection of different atoms in front of the S-L interface. In order to understand the element distribution and segregation of CoCrFeNiCu HEA, EPMA was employed to test the element composition in the selected region of mushy zone and S-L interface (as shown in Fig. 7). The results are shown in Fig. 8, Fig. 9, respectively. Fig. 8 shows the distributions of Cu, Fe, Cr and Co are inhomogeneous in the solidification front. In addition, the distribution of element Ni is uniform in both the dendritic region and inter-dendritic region. It is apparent that phase separation is observed in the solidification front. Cu atoms are piled up in the dendrite tip owing to the positive mixing enthalpy between element Cu and element Co, Cr, Fe during solidification. The content of Cu will reach its solid solubility during solidification. On the other hand, the mixing enthalpy between Ni and Cu is negative. Therefore, the solid solution rich in Cu and Ni will solidify from the liquid CoCrFeNiCu alloy. Similarly, the distribution of Ni in the dendritic regions is approximately equal to that in the inter-dendritic regions, which can be explained by the negative mixing enthalpy between element Ni and other elements. Therefore, Fe, Cr and Co are enriched in the dendritic regions, and Cu, Ni are enriched in the inter-dendritic regions.

Fig. 7.   EPMA selected area of CoCrFeNiCu HEA directionally solidified at 10 μm/s.

Fig. 8.   EPMA-BSE image of the element distribution map in the dendrite tip of CoCrFeNiCu HEA directionally solidified at 10 μm/s.

Fig. 9.   EPMA-BSE image of the elements distribution map of the mushy zone of CoCrFeNiCu HEA directionally solidified at 10 μm/s.

When the sample was quenched into the liquid Ga-In-Sn alloy after directional solidification, some columnar grains with the width of 10 μm were formed in the mushy zone. EPMA results in the mushy zone are shown in Fig. 9, and the element distribution in these columnar grains is similar to that in the dendrite tips as shown in Fig. 8. Cu is enriched in the columnar grain boundaries. The distributions of Cr, Co and Fe in these columnar grains are similar. The distribution of Ni in the columnar grains is the same as that in the grain boundaries. Because of the positive mixing enthalpies between Cu and other elements, Cu atoms would be repelled into the liquid in the mushy zone, and Cu element will be rich in the grain boundaries when the sample was quenched.

According to the above discussion, CoCrFeNiCu HEA contains two phases, CoCrFeNiCu FCC phase and Cu-enriched FCC phase. The phase formation of CoCrFeNiCu HEA can be contributed to two reasons, one is the effect of mixing enthalpy and the other one is the higher diffusion coefficient of element Cu than other elements. The mixing enthalpies between Cu and other elements Co, Cr, Fe and Ni are +6, +12, +13 and +4 kJ/mol, respectively [33]. This means that it is difficult for Cu to thermodynamically mix with other elements, thus resulting in the segregation of Cu [34]. Secondly, the previous literature reports that the diffusion coefficient of Cu is higher than those of other elements, it will promote the precipitation of Cu from liquid alloy during solidification and result in the Cu segregation [35].

Schematic diagram of CoCrFeNiCu HEA solidification process is shown in Fig. 10. The solidified CoCrFeNiCu FCC solid solution coexists with the Cu-enriched liquid alloy in the temperature range of 1394 K―1638 K during solidification. The primary CoCrFeNiCu phase is a simple FCC solid solution structure randomly ordered by all five different elements. As the temperature further decreases, no other phase solidifies from liquid and the primary CoCrFeNiCu FCC solid solution phase continues to grow up to dendritic morphology. Meanwhile, the Cu-enriched liquid alloy is crammed into inter-dendritic region. As the temperature decreases to 1394 K, which is close to the melting point of pure Cu, the Cu-enriched phase will solidify from the liquid alloy. The structure of the Cu-enriched phase is similar to that of pure Cu. The XRD diffraction peaks from other literatures indicate that the Cu-enriched phase mainly includes two types of structures, pure Cu and Cu-rich FCC solid solution [32,34,36,37].

Fig. 10.   Schematic diagram of CoCrFeNiCu HEA solidification process.

In this research, the growing velocity of the primary CoCrFeNiCu FCC phase is significantly different when the solidification velocity increases from 10 μm/s to 100 μm/s, and therefore the precipitation of Cu from the multi-component liquid alloy becomes insufficient. This leads to element Cu fully solubilize into the primary CoCrFeNiCu phase. Meanwhile, the undercooling of the liquid alloy increases obviously with the increasing of solidification velocity, nucleation rate will be increased in front of S-L interface, more nucleation particles will grow into columnar grains with smaller width. As a result, the primary dendritic arm spacing and the secondary dendritic arm spacing decrease obviously when the solidification velocity increased from 10 to 100 μm/s.

3.3. Effect of solidification velocity on tensile properties

Tensile properties of the samples directionally solidified at different velocities were tested, the results are shown in Fig. 11. From this figure, the ultimate tensile strength (UTS) increases from 400 MPa to 450 MPa gradually with increasing velocity from 10 μm/s to 100 μm/s. The elongation increases from 49.5% to 54.7% with the velocity increases from 10 μm/s to 60 μm/s, and then decreases to 50.7% when the velocity is 100 μm/s. It is worth noting that the yield strength of them nearly keeps unchanged, the value is about 290 MPa. The improvements of the UTS and the ductility are mainly resulted from the grain boundary strengthening, more grain boundaries were formed in the refined columnar grains. This strengthening mechanism is different from that of other equiaxed CoCrFeNiCu HEAs strengthened by solid solution and grain boundary. In this research, the effect of solid solution strengthening is weak owing to the serious Cu segregation at grain boundaries. For the grain boundary strengthening, although the number of grain boundaries prepared at higher velocity is more than that prepared at lower velocity, the grain boundaries are parallel to tensile direction, the effect of grain boundary strengthening is greatly weakened, which is accounted for the similar value of the yield strength. The decrease of the strain is caused by the deviation of grain growth direction (from convex or concave S-L interface).

Fig. 11.   Tensile properties of CoCrFeNiCu HEA samples.

Fracture morphologies of different specimens are shown in Fig. 12. There are many dimples with different size on the fracture surface. The fracture morphology of the specimen prepared at 10 μm/s is shown in Fig. 12(a), many tearing ridges and round dimples are uniformly distributed on the fracture surface. This indicates that the fracture mechanism during tensile testing is ductile fracture and cleavage brittle fracture. When the solidification velocity is 30 μm/s, both the number of the dimples and the depth of the dimples increase obviously comparing with those at the 10 μm/s, as shown in Fig. 12(b). As the solidification velocity increases to 60 μm/s, the fracture surface is covered by more dimples and tearing ridges, but the sizes of the dimples and the tearing ridges become smaller. When the solidification velocity further increases to 100 μm/s, the tearing ridges and dimples appear alternatively on the fracture surface (shown in Fig. 12(d)).

Fig. 12.   Fracture morphologies of directionally solidified CoCrFeNiCu HEA specimens: (a) 10 μm/s; (b) 30 μm/s; (c) 60 μm/s; (d) 100 μm/s.

In order to obtain the element distribution of dimple and tearing ridge, energy dispersive spectroscopy (EDS) was used to measure the element distribution on the fracture surface, and the map scanning results of the dimple and the tearing ridge are shown in Fig. 13, Fig. 14, respectively. In Fig. 13, the distributions of element Co, Cr, Fe and Ni are uniform in the selected dimples. However, the concentration distribution of element Cu is higher than other elements. Therefore, it could be concluded that the dimples are originated from the Cu-enriched region during tensile testing. In Fig. 14, the element distribution maps indicate that the contents and distributions of elements in the dimples are significantly different from those on this tearing ridge. The dimple shown on the left side is rich in element Cu, and the tearing ridge shown on the right side is uniformly distributed element Co, Cr, Fe, Ni. To sum up, the dimples are evolved from the Cu-enriched phase and the tearing ridges are formed when the CoCrFeNiCu solid solution phase fractures during tensile testing.

Fig. 13.   EDS result of the fracture surface of directionally solidified CoCrFeNiCu HEA at dimple region.

Fig. 14.   EDS result of the fracture surface of the directionally solidified CoCrFeNiCu HEA at area of the dimple and tearing ridge.

The ductility and strength of the sample solidified at higher solidification velocity are slightly better than those solidified at lower solidification velocity, and the ductility of sample is highest when the sample is directionally solidified at velocity of 60 μm/s. The fracture mechanism of these samples is mixing fracture, which includes ductile and brittle fracture. With the increasing of solidification velocity, the ductile fracture is the dominant mode.

4. Conclusions

CoCrFeNiCu HEA samples were directionally solidified at different solidification velocities. Their microstructures, element distributions, especially the Cu segregation, and tensile mechanical properties were studied. The main conclusions can be drawn as follows:

(1) CoCrFeNiCu HEA can be directionally solidified at these velocities. The S-L interface changes from convex to planar and concave with the increase of solidification velocity, and it is comparatively planar at velocity of 30 and 60 μm/s. The morphology of the S-L interface is mainly influenced by axial heat transfer and solidification velocity during directional solidification.

(2) The primary and the secondary dendritic spacings decrease from 100 μm to 10 μm and from 20 μm to 5 μm respectively, which is caused by grain competition growth during directional solidification.

(3) Element Cu is enriched in the inter-dendritic region and solidification front zone, and Cu content increases with the solidification velocity. Element Cu is compelled from the FCC phase and accumulates in the liquid owing to its low mixing enthalpy during directional solidification.

(4) UTS increases gradually from 400 MPa to 450 MPa with the increase of solidification velocity, and the strain of the specimens with the velocity of 60 μm/s is higher than that of others.

(5) The improvement of UTS results from columnar grain boundary strengthening. The decreases of the strain for the samples prepared at 10 and 100 μm/s are caused by the deviation of grain growth to pulling direction.

Acknowledgments

This work was supported financially by the National Natural Science Foundation of China (Nos. 51825401and51741404) and the State Key Laboratory of Advanced Welding and Joining.


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