Microstructure and mechanical properties of ultra-fine grained MoNbTaTiV refractory high-entropy alloy fabricated by spark plasma sintering

1. Introduction

The refractory high-entropy alloy (RHEA) was proposed by Senkov et al. [1] based on the conception of high-entropy alloy (HEA), which is consisted of the refractory elements, such as Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W [2]. Compared with conventional high-temperature alloys, the RHEAs exhibit higher melting point, excellent high temperature mechanical properties [1,3], outstanding wear, oxidation and corrosion resistance [[4], [5], [6]]. MoNbTaW and MoNbTaVW RHEAs maintained yield strengths of 405 MPa and 477 MPa at 1600 °C respectively, which were noticeable superior than other high-temperature alloys [3]. Compared with commercial superalloy (Inconel 718), the MoTaWNbV RHEAs fabricated by Poulia et al. [7] exhibited greater wear resistance. The NbCrMo_0.5Ta_0.5TiZr RHEA performed a better oxidation resistance than those of commercial Nb alloys, NbSiAlTi and NbSiMo alloys under the condition of 1000 °C for 100 h in flowing air [4]. These advantages suggest significant potential for high-temperature applications, such as the aerospace high-temperature load-bearing structures and thermal protection structures. To date, the fabrication methods of the RHEAs are mainly focused on vacuum arc-melting due to the high melting points of component elements [[8], [9], [10]]. However, the segregation, dendritic microstructures and coarse grains are hardly controlled in most of arc-melting RHEAs. These undesirable microstructures generally deteriorated the mechanical properties of alloys, which limits the potential applications of the RHEAs.

Powder metallurgy (PM) is an effective method to fabricate bulk alloys [[11], [12], [13]]. The ultra-fine, or even nanocrystalline alloys and solid solution powders can be synthesized by mechanical alloying (MA) using metal powders. Combined with spark plasma sintering (SPS), the mechanically alloyed powders can be rapidly consolidated in a few minutes [[14], [15], [16]]. The submicron scale grains and homogeneous microstructures are obtained, which effectively enhances the mechanical properties of alloys [[17], [18], [19]]. Based on these advantages of MA and SPS, Kim et al. [20] reported a CoCrFeMnNi high-entropy alloy. The bulk alloys consolidated at 900 °C using 60 h MA powders owned an average grain size of 270 nm. The compressive fracture strength arrived 3 GPa which was three times higher than that fabricated by melting method. Moravcik et al. [21] successfully synthesized ultra-fine grained Ni_1.5Co_1.5CrFeTi_0.5 HEA using MA and SPS. The alloy showed excellent mechanical properties, such as a bend strength of 2593 MPa, a tensile strength of 1384 MPa and an elastics modulus of 216 GPa. Therefore, it is expectable to fabricate RHEAs with ultra-fine grains and homogenous microstructures using MA and SPS to enhance mechanical properties.

In this work, based on the conventional vacuum arc-melting MoNbTaTiV RHEA with promising mechanical properties [9], the MA and SPS were used with an aim to obtain ultra-fine grains and homogeneous microstructures, as well as further enhanced mechanical properties. The phase evolution, microstructural characteristics and mechanical properties at room temperature of the MoNbTaTiV RHEA were systematically investigated. The strengthening mechanisms were also discussed and revealed.

2. Experimental

The MoNbTaTiV RHEA powders were fabricated by high-energy ball milling using the elemental powders of Mo, Nb, Ta, Ti and V. The characteristics of these original powders are shown in Table 1. The component elemental powders were placed in a planetary ball miller (QM-3SP4, Nanjing NanDa Instrument Plant) with stainless steel vials and balls. The MA process was carried out with a ball-to-powder weight ratio of 10:1 at 350 rpm for 40 h. The argon atmosphere was used to protect alloys from oxidation. The miller was interrupted per 30 min and halted for another 15 min to avoid overheating. No process control agent (PCA) was used to ensure the purity of the mechanically alloyed powders.

The mechanically alloyed powders were subsequently sintered using SPS in a graphite mold with an inner-diameter of 15 mm. The SPS processes were carried out at 1500 °C, 1600 °C and 1700 °C for 10 min with an axial pressure of 50 MPa under vacuum (< 1.3 × 10^-2 Pa). The mechanically alloyed powders were heated to 700 °C within 3 min, then the heating rate is 100 °C/min from 700 °C to 1500 °C, 1600 °C and 1700 °C. The sintering temperature during SPS was measured by an infrared pyrometer.

The crystal structures of the mechanically alloyed powders and bulk alloys were characterized by X-ray diffraction (XRD) using a D/max-rb diffractometer with Cu Kα radiation at 40 kV and 40 mA in the 2θ range from 20° to 80°. The microstructures of the mechanically alloyed powders and bulk alloys were observed by a Zeiss supra55 scanning electron microscopy (SEM), while the chemical composition was measured by an energy dispersive spectrometer (EDS). Transmission electron microscopy (TEM) analysis was carried out by a Talos F200x operated at 200 kV to observe microstructures at nanoscale. The impurity elements of O and N in the bulk alloys were determined by an ONH836 Nitrogen/Oxygen Determinator. The densities of the bulk alloys were determined using Archimedes method. The HVS-1000A Vickers micro-hardness tester was used to determine the hardness of the bulk alloys at a load of 9.8 N with a loading time of 15 s. The compressive tests were carried out at room temperature using an Instron 5569 universal testing machine with cylindrical samples (2.5 mm diameter and 4 mm height) at an engineering strain rate of 10^-3 s^-1. The densities, micro-hardnesses and compressive tests were conducted three times at each condition to obtain the average values.

3. Results and discussion

3.1. Phases and characteristics of MoNbTaTiV RHEA powders

Fig. 1(a) shows the XRD patterns of the MoNbTaTiV RHEA powders. The component elements can be observed from the blended powders. The mechanically alloyed powders with simple BCC phase are obtained after 40 h milling process. The drastically declined and broadened of diffraction peaks are mainly attributed to the refinement of the grain size and the increment of the lattice strain. According to the Williamson-Hall method, the grain size and the lattice strain are estimated to be 12.9 nm and 0.89, respectively. The nano-scale grain size and severe lattice strain lead to the mechanically alloyed powders into hardening which results in the drastically refined of powders. As shown in Fig. 1(b), the mechanically alloyed powders with an average size of 6.8 μm are consisted of a large number of refined powders and relatively few larger powders. These refined powders can be effectively filled in the gap between the larger powders. The approximate spherical morphologies of the mechanically alloyed powders indicate great fluidity. These factors are beneficial to the consolidation quality. Furthermore, the repeated cold-welding and fracture greatly promote the diffusion and dissolving of the component elements. The EDS analysis of the mechanically alloyed powders is listed in Table 2. Each component element has almost same level compared with the nominal composition, which confirms relatively uniform elemental distribution.

3.2. Phases and microstructures of the bulk MoNbTaTiV RHEAs

After SPS process, the bulk alloys still exhibit simple BCC phase indicating the satisfactory thermal stability of the mechanically alloyed powders (Fig. 2). The remarkable sharp and high intensity diffraction peaks are mainly associated with the growth of the grain size and the release of the lattice strain. The extremely small FCC diffraction peaks which are closely with TiO and TiN are detected from the bulk alloys sintered at 1700 °C.

As shown in Fig. 3, the bulk alloys are consisted of the gray matrix phase and uniformly dispersed dark precipitation phases. The extremely refined precipitation phases decline the intensities of the XRD diffraction peaks, which results in the unsatisfactory observation of XRD for the bulk alloys sintered at 1500 °C and 1600 °C. The precipitation phases are gradually coarsened with the increasing of sintering temperature. Until the extremely small diffraction peaks are detected from the bulk alloys sintered at 1700 °C. Furthermore, the volume fractions of precipitation phases are decreased from 4.94% to 3.88% as the sintering temperature increases from 1500 °C to 1700 °C. This trend is attributed to the dissolution of precipitation phases at high temperature.

The TEM bright field images and corresponding selected area electron diffraction (SAED) patterns of the MoNbTaTiV RHEAs sintered at different temperatures are shown in Fig. 4. The precipitation phases are dispersed in both boundaries (marked in green circle) and interior (marked in red circle) of the submicron matrix grains. As the sintering temperature increases, the grains of both the matrix and precipitation phases are gradual growth. The lattice parameter of the MoNbTaTiV RHEA is calculated to be 3.218 Å using the ROM [9]. According to the SAED pattern, the lattice parameter of the matrix is calculated to be 3.206 Å, which is closely to the theoretical value. The SAED pattern of the precipitation phase reveals FCC structure with the lattice parameter of 4.188 Å. Compared with the PDF card of TiO (4.177 Å) and TiN (4.24 Å), the relative moderate value probably indicates the TiO and TiN compounds.

The component elements of the bulk alloys are determined by high angle annular dark field (HAADF) image and corresponding EDS maps in Fig. 5. The Mo, Nb, Ta, V and relatively few Ti, O and N are uniformly distributed in the matrix. More Ti, O and N are concentrated in the regions of precipitation phases. Thus, it is confirmed that the precipitation phases are the TiO and TiN compounds. Actually, the active Ti possesses extremely affinity towards air. During MA, the reaction between Ti and residual air is promoted by repeated fracture and cold-welding to form the compounds, which are gradually refined, dispersed and dissolved into the matrix, then precipitated during SPS. Most of the O and N are distributed in the compounds. However, the high temperature, concentration defects and stored energy also promote the diffusion of the O and N into the matrix [22,23].

The average grain sizes and volume fractions of the matrix and precipitation phases are counted and listed in Table 3, as well as the contents of O and N. The ultra-fine grain sizes of the bulk alloys are mainly attributed to the nanocrystalline mechanically alloyed powders by the MA process and the rapid consolidation of the SPS process. Furthermore, the extremely refined and uniformly dispersed compounds are also beneficial to inhibit the rapid growth of grain size at high sintering temperatures [24].

3.3. Mechanical properties of bulk MoNbTaTiV RHEAs

After SPS process, the bulk alloys sintered at different temperatures exhibit excellent mechanical properties. The maximum micro-hardness reaches 542 HV (Fig. 6). Compared with vacuum arc-melting method, the micro-hardness value is enhanced 22.35% [9]. The relative small error bounds are attributed to the homogeneous microstructures.

The compressive engineering stress-strain curves of the MoNbTaTiV RHEAs at room temperature are shown in Fig. 7. The corresponding yield strengths, fracture strengths, plastic strains and densities are listed in Table 4. The bulk alloys present outstanding yield and fracture strengths, as well as the acceptable compressive ductility. The distinct stress overshoot is observed in the stress-strain curves. This phenomenon is similar to that observed also for many steels with solute carbon, and results here may from the interaction between interstitial atoms and the dislocations [25]. Different from the vacuum arc-melting method, the diffusion of atom is probably suppressed due to the lower processing temperature. The limited consolidation time of the SPS restricts the consolidation process. The inherent sluggish diffusion effect of HEAs can reduce the diffusion coefficient and increase activation energy [26]. Thus, these factors lead to the insufficient consolidation at 1500 °C, which results in the decrease of density and mechanical properties. At the end of the compressive engineering stress-strain curves, the obvious fracture steps are observed. The crack along the direction of compression occurs at internal of the bulk alloys at the first step, which is considered as the fracture strength. With the generation of crack, the stress is slightly decreased. However, the integrality of the bulk alloys are not affected by the internal crack, which leads to the slightly increase of compressive stress until the generation of other cracks. Eventually, with the addition of cracks, the integrality of the bulk alloys are destroyed, which results in the rapid decrease of compressive stress.

The corresponding compressive fracture morphologies of the MoNbTaTiV RHEAs at room temperature are shown in Fig. 8. The typical flabelliform river patterns are observed in all bulk alloys, indicating the cleavage fracture which generally happened in the alloys with BCC structure. The apparent intergranular fracture morphologies are observed in the corresponding magnified images (Figs. 8(d)-(f)). The ultra-fine grains are accordance with the results of TEM images. A large number of compounds particles are uniformly dispersed in the grain boundaries of the matrix. Different from the compact compressive fracture morphologies of the bulk alloys sintered at 1600 °C and 1700 °C, the evident micro-cracks are observed from the magnified images of the bulk alloys sintered at 1500 °C, which is attributed to the lower binding force by the insufficient consolidation.

The room temperature mechanical properties of some typical RHEAs and MoNbTaTiV RHEA sintered at 1600 °C are summarized in Table 5. The corresponding compressive yield strengths and plastic strains are depicted in Fig. 9. It is evident that the RHEAs fabricated by MA and SPS perform better mechanical properties compared with those fabricated by arc-melting method. The yield and fracture strengths of the MoNbTaTiV RHEA are respectively enhanced 57.5% and 32.2% than the arc-melting processed [9]. However, the large number of uniformly dispersed and refined compounds destroy the continuity of matrix, which results in the slightly decrease of plastic strain.

As reported, the outstanding mechanical properties of the RHEAs fabricated by MA and SPS are mainly related to several strengthening mechanisms: the substitution solid solution strengthening attributed to the inherent concept of HEAs, the precipitation strengthening by the precipitated compounds, the grain boundary strengthening followed by the Hall-Petch relationship and the interstitial solid solution strengthening by the interstitial atoms [23]. All of these strengthening mechanisms are independent. The yield strength of the MoNbTaTiV RHEA can be estimated as the summation of each strengthening mechanism and is expressed as:

σ_y=σ₀+Δσ_ss+Δσ_or+Δσ_gb+Δσ_iss (1)

where σ₀ is the intrinsic strength related to the lattice frictional strength, and Δσ_ss, Δσ_or, Δσ_gb and Δσ_iss are the contributions of the substitution solid solution strengthening by component elements, the precipitation strengthening by compounds, the grain boundary strengthening by ultra-fine grain and the interstitial solid solution strengthening by interstitial elements, respectively.

The contribution of precipitation strengthening mechanism is more suitable estimated using Orowan dislocation bypass mechanism because of the high shear modulus and hardness of the compounds. The Orowan strengthening mechanism is described by the following equation [32]:

σ_or=$M\frac{0.4Gb}{πλ}⋅\frac{ln(2\bar{r}/b)}{(1-ν)^{1/2}}$ (2)

where M (2.9) is the average orientation factor for the BCC polycrystalline matrix [33], G is the shear modulus, ν is the Possion's ratio, and b is the Burgers vector of the matrix. Considering that some parameters for the MoNbTaTiV RHEA are not directly obtained from the literatures. The values of G (75.1 GPa) and ν (0.361) are approximately used according to the literature [34]. Burgers vector b is equal to (2/3)^1/2a, and a is the lattice parameter for a BCC structure. -r is the average radius of a circular cross-section in a random plane for a spherical precipitate, and is equal to (2/3)^1/2r. λ is the interspacing of a precipitate, and is calculated as follows [35]:

λ=$2\bar{r} ((\frac{π}{4f})^{1/2}-1)$ (3)

where f is the average volume fraction of the precipitation phase. Therefore, according to the Eqs. (2) and (3), the calculated values of the σ_or are 119 MPa, 91 MPa and 51 MPa for the bulk alloys sintered at 1500 °C, 1600 °C and 1700 °C.

The interstitial solid solution strengthening is proportional to the square root of interstitial atoms contents according to the well-accepted Fleischer model [22,23,25]. However, considering the complexity of the microstructures and the uncertainty of the corresponding parameters in the Fleischer model, the simplified equation is also acceptable for the HEA, and can be approximately expressed as [22,23]:

Δσ_iss=$Qc_m^{1/2}$ (4)

where Q is a material constant, c_m is the atomic concentration of interstitial atoms in the matrix. Combined with the results of the chemical analysis and volume fractions of the precipitation phases, the contents of O and N in the matrix are approximately calculated and listed in Table 6. The gradual increased interstitial atoms in the matrix are probably attributed to the promotion of the higher temperature. Actually, the interstitial solid solution strengthening is mainly due to the lattice distortion generated by the interstitial atoms and the matrix, which hinters the movement of dislocations. The lattice distortion is related to the radius of interstitial atoms [36]. In this work, both of the O and N atoms are diffused into the matrix, which is difficult to accurately determine the content of each element. To simplify the analysis, we assume that O and N are as one interstitial element in the interstitial solid solution strengthening mechanism, considering their adjacent atomic number and similar characteristic, such as atomic radius, valence electron concentration and electronegativity [26].

Generally, the values of σ₀ and σ_ss are considered as intrinsic for a certain HEA, which are almost independent of the fabrication methods. Therefore, after excluding σ_or, σ_iss, σ₀ and σ_ss, the difference of yield strength is only ascribed to the grain boundary strengthening. The values can be estimated by the Hall-Petch formula as [37]:

σ_gb=k_yd^-1/2 (5)

where k_y is the Hall-Petch coefficient and d is the average grain size.

In this work, in order to calculate the values of σ₀, σ_ss, σ_iss, and σ_gb of the MoNbTaTiV RHEA, we assume that no Orowan strengthening and interstitial solid solution strengthening in the arc-melting MoNbTaTiV RHEA which exhibits an average grain size about 100 μm [9]. Therefore, according to the Eqs. (1)-(5), the Q is calculated to be 432 MPa. The relationship between the compressive yield strength and grain size of the bulk alloys is illustrated in Fig. 10. The value of the line intercept is 1376 MPa which indicates the combination of the intrinsic strength and substitution solid solution strength. The slope of the line is 241 MPa μm^0.5 which indicates the Hall-Petch coefficient. It should be noted that the bulk alloys sintered at 1500 °C are useless because of the insufficient consolidation. The contributions of these strengthening mechanisms are summarized in Table 6.

It is clear that the grain boundary strengthening is the inherent advantage of the MA and SPS, which effectively enhances the yield strength for BCC structure alloys [38]. The interstitial solid solution strengthening also plays an important role in the MoNbTaTiV RHEA. Actually, the lower diffusion activation energy and the smaller atomic packing factor of BCC structure greatly promote the diffusion of interstitial atom, which leads to the enhancement of yield strength. In addition, the samples sintered at 1700 °C with significantly coarsened grains have similar yield strength to those sintered at 1600 °C, after excluding the contribution of precipitates. Based on the content of interstitial atoms in the matrix (Table 6), it should be attributed to the stronger interstitial solid solution strengthening in the samples sintered at 1700 °C compensated the weakening of grain boundary strengthening. This supports the assumptions and calculations of the strengthening mechanisms. Therefore, it can be concluded that the interstitial solid solution strengthening and grain boundary strengthening are the main strengthening mechanisms in the MoNbTaTiV RHEA fabricated by MA and SPS.

4. Conclusions

(1) The MoNbTaTiV RHEA with ultra-fine grains and homogenous microstructures was successfully fabricated by MA and SPS. The refined mechanically alloyed powders with nanocrystalline and simple BCC phase were obtained after 40 h MA process. The microstructures of the bulk alloys were consisted of submicron scale BCC matrix and FCC precipitation phases which were uniformly dispersed in both boundaries and interior of the matrix grains. The precipitation phases were mainly consisted of TiO and TiN. These submicron scale grains and homogeneous microstructures were attributed to the advantages of MA and SPS.

(2) The MoNbTaTiV RHEA fabricated by MA and SPS exhibited outstanding mechanical properties. The micro-hardness, compressive yield strength, fracture strength and plastic strain of the bulk alloys sintered at 1600 °C at room temperature are 542 HV, 2208 MPa, 3238 MPa and 24.9%, respectively. Compared with vacuum arc-melting method, the enhanced of mechanical properties of the MoNbTaTiV RHEA fabricated by MA and SPS were mainly attributed to the grain boundary strengthening by ultra-fine grains and the interstitial solid solution strengthening by interstitial elements.

Acknowledgements

This work was supported financially by the National Natural Science Foundation of China (No. 51875122).

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