III. RESULTS AND DISCUSSION

Figure 1(a) shows the XRD pattern of the as-cast CrMoNbTaV HEA. According to the XRD pattern analysis, it reveals four sharp diffracted peaks at 2θ = 39.7°, 57.5°, 72.2°, and 85.2°, which corresponded to the reflection planes of (110), (200), (211), and (220), respectively. This observation confirms that the as-cast alloy has a single BCC structure like many other refractory HEAs such as, HfNbTiVSi_0.5,^{Reference Zhang, Liu, Li, Chen and Zhang21} TiZrHfNbTa,^{Reference Couzinié, Lilensten, Champion, Dirras, Perrière and Guillot23} Mo_0.5NbHf_0.5ZrTi,^{Reference Guo, Wang, Luo, Li, Chen, Su, Guo and Fu24} and MoNbHfZrTi.^{Reference Guo, Wang, Luo, Li, Chen, Su, Guo and Fu25} And the lattice constant of the BCC structure is 0.320 nm. Back scattered-electron (BSE) SEM images of the as-cast alloy are presented in Figs. 1(b) and 1(c). It can be seen that the dendritic microstructure is obviously formed in the CrMoNbTaV alloy, which is typical of as-cast ingots. The dendrite arms appear to be more brighter (in terms of SEM image contrast) than the interdendritic regions. The dendrites are slightly rounded and rather branched in shape. The EDS analysis results show that the dendrite region is enriched in Mo and Ta elements with relatively higher melting points. By contrast, the interdendritic region is enriched in Cr and V elements which have lower melting points. The distribution of the Nb element with moderate melting point between dendrite and interdendritic is relatively uniform.

FIG. 1.

XRD pattern (a) and BSE SEM images of the as-cast HEA CrMoNbTaV: (b) low magnification SEM image and (c) high magnification SEM image with EDS analysis results (at.%).

Based on the previous studies,^{Reference Guo, Hu, Ng and Liu26–Reference Singh and Subramaniam29} there are five parameters in formulating thermodynamic rules for the formation of the HEA phase: the atomic radius difference (δ), the mixing enthalpy (ΔH _mix), the mixing entropy (ΔS _mix), the electronegativity difference (Δχ), and the valence electron concentration (VEC). They are defined by Eqs. (1)–(5), respectively.

(1)$$\delta = 100\sqrt {\sum\limits_{i = 1}^n {{c_i}{{\left( {1 - {{{r_i}} / {\bar{r}}}} \right)}^2}} } ,\quad \bar{r} = \sum\limits_{i = 1}^n {{c_i}{r_i}} \quad ,$$

where, c _i is the at.% of the ith element, r _i is the atomic radius of the ith element, $\bar{r}$ is the average atomic radius of all the alloying elements in the present system, and n is the number of components. The numerical factor 100 is used to amplify the data for clarity; and

(2)$$\Delta {H_{{\rm{mix}}}} = \sum\limits_{i = 1,i \ne j}^n {{\Omega _{ij}}{c_i}{c_j}} ,\quad {\Omega _{ij}} = 4\Delta H_{{\rm{mix}}}^{{\rm{AB}}}\quad ,$$

where $\Delta H_{{\rm{mix}}}^{{\rm{AB}}}$ is the mixing enthalpy of binary AB atomic pairs; and

(3)$$\Delta {S_{{\rm{mix}}}} = - R\sum\limits_{i = 1}^n {{c_i}\ln {c_i}} \quad ,$$

where R is the universal gas constant; and

(4)$$\Delta \chi = \sqrt {\sum\limits_{i = 1}^n {{c_i}{{\left( {{\chi _i} - \bar{\chi }} \right)}^2}} } ,\quad \bar{\chi } = \sum\limits_{i = 1}^n {{c_i}{\chi _i}} \quad ,$$

where χ_i is the Pauling electronegativity of ith element and $\bar{\chi }$ is the average Pauling electronegativity of all the alloying elements in the present system; and

(5)$${\rm{VEC}} = \sum\limits_{i = 1}^n {{c_i}{{\left( {{\rm{VEC}}} \right)}_i}} \quad ,$$

where (VEC)_i is the valence electron concentration of ith element.

The at.% (c _i), the atomic radius (r _i), the Pauling electronegativity (χ_i), and the valence electron concentration [(VEC)_i] of each element are exhibited in Table I. Meanwhile, as shown in Table II, the chemical mixing enthalpy $\left( {\Delta H_{{\rm{mix}}}^{{\rm{AB}}}} \right)$ of the binary equiatomic alloy was calculated based on Miedema’s model proposed by Sun et al.,^{Reference Sun, Yi, Liu, Zang and Jiang30} to calculate the thermodynamic parameters for the present system. According to Eqs. (1)–(5), the thermodynamic parameters of the CrMoNbTaV HEA are calculated in Table III.

TABLE I.

The at.%, the atomic radius, the Pauling electronegativity, and the VEC of each element in HEA CrMoNbTaV.

TABLE II.

The chemical mixing enthalpy $\left( {\Delta H_{{\rm{mix}}}^{{\rm{AB}}}} \right)$ of binary equiatomic alloys calculated by Miedema’s approach.

TABLE III.

Calculated parameters δ, ΔH _mix, ΔS _mix, Δχ, and VEC for CrMoNbTaV HEA.

Guo et al.^{Reference Guo and Liu27} showed that solid solution phases only form when the three parameters, δ, ∆H _mix, and ∆S _mix, are satisfying these criteria: 0 ≤ δ ≤ 8.5, −22 ≤ ∆H _mix ≤ 7 kJ/mol and 11 ≤ ∆S _mix ≤ 19.5 J/(K mol), respectively. In this study, the calculated values of δ (6.95), ∆H _mix (−4.64 kJ/mol), and ∆S _mix [13.38 J/(K mol)] are well within the criteria. Guo et al.^{Reference Guo and Liu27} also found that the valence electron concentration (VEC) plays a decisive role in determining the FCC- or BCC-type solid solution in HEAs, and specially larger VEC (≥8) favors the formation of FCC-type solid solutions, while smaller VEC (≤6.87) favors the formation of BCC-type solid solutions. FCC and BCC coexist when VEC is larger than 6.87 and less than 8. The value of VEC is 5.4 in the present work, which indicates that BCC solid solution could exist steadily in our synthesized alloy, agreeing with the experimental results.

However, Guo et al.^{Reference Guo and Liu27} thought that Δχ almost does not have an effect on the formation of the solid solution phase. But Singh et al.^{Reference Singh and Subramaniam29} found that Δχ is also an important factor in predicting solid solution formation; when Δχ ≥ 10%, compound phases are also likely to form along with HEA solid solution. In this study, Δχ is calculated to be 23.13%; however, reflections of BCC single solid solution without compound phase are identified for the as-cast CrMoNbTaV HEA [Fig. 1(a)]. This result is well consistent with the criteria for the formation of solid solution of Guo et al.^{Reference Guo and Liu27}

Microhardness measurements should provide first insight into the mechanical properties of the alloy at room temperature,^{Reference Gorr, Azim a, Christ, Mueller, Schliephake and Heilmaier31} carried out on the as-cast sample. As expected, the interdendric regions show higher microhardness values compared to the dendrite areas. As shown in Fig. 2, the average microhardness value of the interdendritic regions is about 950 HV, while the corresponding value for the dendrite regions is about 896 HV. It may be caused by the content of element Cr, which in the interdendritic regions is higher than that in the dendrites regions. Compared to other HEAs such as TaNbHfZrTi (390 HV),^{Reference Senkov, Scott, Senkova, Miracle and Woodward16} AlCrFeNiMo_0.2 (465 HV),^{Reference Dong, Jiang, Jiang, Lu, Wang and Li32} and NbCrMo_0.5Ta_0.5TiZr (539 HV),^{Reference Senkov and Woodward33} the microhardness value of this alloy is much higher for the reason of the large difference of the atomic radius between the element Cr and other four elements. This difference makes the alloy has more significant lattice distortion effect, which strengthens the alloy so as to improve the microhardness. The microhardness indentation carried out on the as-cast sample clearly demonstrates the ductile nature of the alloy since cracks are neither detected in the dendrites nor in the interdendritic regions.

FIG. 2.

The microhardness values and SEM images of dendrite and interdendrite regions of the as-cast HEA CrMoNbTaV.

Figure 3 shows thermogravimetric curves of the CrMoNbTaV HEA in the as-cast condition during exposure to laboratory air at 900, 1000, and 1100 °C. The oxidation kinetic curves at 900 and 1000 °C show the same tendency that the mass gain increases as the exposure time extends. The oxidation kinetic curves obtained at 900 and 1000 °C obey pseudo-parabolic rate law indicating that the growth of the oxide scale proceeds through solid state diffusion. And the oxidation rate at 1000 °C is substantially higher than that at 900 °C, and it reveals that the oxide scales belong to the nonprotective oxide scales.

FIG. 3.

Oxidation kinetics curves of the HEA CrMoNbTaV.

To understand whether a simple kinetics law can apply, the curves can be fitted by an equation model^{Reference Liu and Xue34} as

(6)$$\Delta {M_n} = {k_n}t\quad ,$$

where ΔM _n is the mass grain per unit surface area of the specimen, t is the exposure time, and k _n is the rate constant. The rate constant k _n and exponent n were evaluated from the linear regression fitting of ln(ΔM) versus ln t data. And the relationship^{Reference Liu and Xue34} of the two data is as follows:

(7)$$\ln \left( {\Delta M} \right) = {1 \over n}\ln t + {1 \over n}\ln {k_n}\quad .$$

And after taking logarithm of mass gain and exposure time, the oxidation kinetic curves of this alloy at 900 and 1000 °C obey pseudo-parabolic growth kinetics, and the values of n of this alloy after oxidation at 900 and 1000 °C for 25 h are 0.95 and 1.00, respectively. However, the thermogravimetric curve for 1100 °C shows a different tendency compared to that at 900 and 1000 °C. With the exposure time increasing, the mass change is found to increase first at 5 h, followed by continuous reduction. This may result from the fact that the oxide scale of this alloy after 1100 °C oxidation is likely to be detached from the matrix with the oxidation time increasing. It is proved that the oxidation behavior of this alloy is greatly influenced by the oxidation temperature.

Figure 4 presents the surface morphologies of the oxides after oxidation at 900, 1000, and 1100 °C for 25 h. The SEM micrographs reveal that a thick and continuous oxide-rich region is formed on the surface of this alloy, and the morphologies of the oxides formed at 900, 1000, and 1100 °C are similar. After oxidation at these three temperatures, long rod-shaped oxidation products are observed on alloy surface. It can be seen that the size of long rod-shaped oxidation products at 1100 °C is somewhat larger than those of 900 and 1000 °C. One can note that all the oxidation products at these three temperatures show a trend of directional growth. The different oxides that are observed by the energy spectra analysis are marked in Figs. 4(b), 4(d), and 4(f). The EDS results show that most of the oxides are rich in Ta, Nb, and V. With the oxidation temperature increasing, the content of O increases gradually while the contents of Nb, Ta, and V decrease.

FIG. 4.

BSE SEM images and EDS analysis results (at.%) of the surface morphologies after oxidation at (a and b) 900 °C, (c and d) 1000 °C, (e and f) 1100 °C.

Figure 5 shows the XRD results of the surface oxides of this alloy after oxidation at 900, 1000, and 1100 °C for 25 h. It can be noted that the oxidation products formed on the alloy surfaces are similar at 900 and 1000 °C, being mainly Nb₂O₅, NbO₂, CrTaO₄, CrNbO₄, Ta₉VO₂₅, Nb₉VO₂₅, and TaO. As the oxidation temperature is elevated to 1100 °C, it can be seen clearly that all the peaks of CrTaO₄ and CrNbO₄ disappeared. The disappearance of CrTaO₄ and CrNbO₄ in the oxidation products at 1100 °C can be mainly attributed to the detachment of oxide scale from the matrix caused by oxidation temperature increasing.

FIG. 5.

XRD pattern of the CrMoNbTaV HEA after oxidation for 25 h.

To further investigate the oxide-rich region, the characteristics of cross sections are analyzed by SEM in Fig. 6 with corresponding EDS in Table IV. The images in Fig. 6 reveal that a thick and continuous oxide scale is formed on the surface of this alloy after oxidation at 900, 1000, and 1100 °C for 25 h. The thicknesses of oxide scales are measured to be about 530 μm, 1132 μm, and 600 μm at 900 °C, 1000 °C, and 1100 °C, respectively. From the pictures we can see that the morphologies of oxide scales at different oxidation temperatures are very similar. All the oxide scales can be divided into three regions: (I) a thin and attached to the substrate region, (II) a thick and porous oxide mixture, and (III) a relatively thin outer region. The innermost layer (region I), usually a diffusion layer, is thin and dense. A distinct crack is observed in the region I after oxidation at 900 and 1000 °C, while more cracks are formed in it as the oxidation temperature increased to 1100 °C. Region II with the largest thickness is composed of two microstructures, namely gray and white regions. In region II, more pores can be observed as compared with the case in regions I and III, indicating that the structure of oxide scale is looser. The outer oxide layer (region III) consists almost entirely of white phase. In addition, all the oxide scales at these three temperatures exhibit a slight tendency to flake.

FIG. 6.

BSE cross-sectional micrographs of the scales formed on the CrMoNbTaV HEA oxidized for 25 h at (a) 900 °C, (b) 1000 °C, and (c) at 1100 °C.

TABLE IV.

Quantitative EDS chemical analyses (at.%) of the oxidation products indicated in Fig. 6.

The EDS analysis of the oxide scales in Table IV reveals that the scales consist of a mixture of Cr, Mo, Nb, Ta, and V oxides. The enrichment of Mo is only found at the interface oxide/substrate, which indicates that the evaporation of Mo oxides can be neglected.^{Reference Gorr, Mueller, Christ, Mueller, Chen, Kauffmann and Heilmaier9} According to the EDS analysis, the gray region in region II consists of a mixture of Cr, Nb, Ta, and V oxides, and the white region consists of a mixture of Nb, Ta, and V oxides. To sum up, the oxide scale consists of three regions: (I) a mixture of Cr, Mo, Nb, Ta, and V oxides, (II) a mixture of Cr, Nb, Ta, and V oxides, and (III) a mixture of Nb, Ta, and V oxides.

Based on the above results and discussion, the oxidation behavior of the CrMoNbTaV HEA can be summarized as follows. There is a thick and porous oxide layer formed on the surface of this alloy after oxidation at high temperature ranging from 900 °C to 1100 °C. This is mainly because that this alloy contains Mo and V elements, both of which have negative effects on the high-temperature oxidation behavior of the metallic material. The chemical stability of the Mo element is relatively lower, and Mo will react with O in the air when the temperature increases to 600 °C. The corresponding oxidation product MoO₃ is volatile with low melting point of 795 °C. The volatile process of MoO₃ can cause the structure of oxide scale loose and porous so that the alloy surface can not form a dense oxide thin layer. As a result, the alloy will have been “destructive” oxidation.^{Reference Li, Zhang, Zhao and Jia35} This is also the reason why there is no Mo element observed on the surface of the oxide layer. Adding too much V element will be adverse to the high temperature oxidation resistance of the alloy. This is because that V oxides including VO, VO₂, and V₂O₅ all have a low melting point and easy to volatilize at high temperature. With the oxidation temperature increasing, V oxides come out from the alloy oxide scale, which destroys the formation of the dense oxide layer.^{Reference Lee36}