1 Introduction

Biomedical materials are extensively used as surgical instruments, cardiovascular stents, and dental and orthopedic implants due to their good processing properties and superior mechanical properties [1,2,3]. Despite considerable success in biomedical applications, the current biomedical metallic materials (e.g., Co–Cr alloys, Ni–Ti alloys, Ti alloys, and stainless steels) still suffer from some problems, such as inflammatory responses or Alzheimer’s disease that may be induced by the release toxic ions of Al and V, particle disease resulted from corrosion and friction of the alloys, stress-shielding effect caused by the mismatch of Young's moduli between the implants and bones [2,3,4]. It is of importance to develop new metallic biomaterials with high strength, low elastic modulus, good corrosion resistance, and biocompatibility for fulfilling the demand of biomedical applications.

High-entropy alloys (HEAs), which are defined as the alloys containing five or more principal elements, have attracted much attention due to their high strengths and hardness, good ductility, and superior wear and anti-corrosion properties [5,6,7,8,9]. From the perspective of biocompatibility, the HEAs composed of biocompatible elements, such as Ta, Nb, Mo, Zr, Ti, and Hf, are attractive [10,11,12,13,14,15,16,17,18,19]. In addition, these elements are the usual constituent elements of the body-centered-cubic (BCC) HEAs. In the last decade, a series of HEAs in Ti–Zr–Hf–Nb–Ta [11, 20,21,22,23], Ti–Mo–Ta–Nb–Zr [18, 19, 24, 25], and Ti–Nb–Hf–Ta–Zr–Mo [26, 27] systems with suitable mechanical and chemical properties for biomedical applications have been studied. Efforts have also been devoted to improving the properties of the HEAs by alloying with O, Si, Al, and Cr [7, 10, 28,29,30,31,32,33]. It is noticed that the addition of Sn in the Fe–Co–Cu–Ni(–Mn) HEAs enhanced the tensile strengths [34,35,36]. Meanwhile, Sn is non-cytotoxic and commonly involved in the β-Ti alloys, which are used as biomedical materials with relatively low Young’s moduli and good corrosion resistance and biocompatibility [37,38,39,40,41,42,43,44]. Recently, we have found that the equimolar TiZrHfNbTa HEA exhibited low elastic modulus, good corrosion resistance, and in vitro biocompatibility [23]. In the present work, with the aim of synthesizing novel biomedical alloys with the integration of improved mechanical properties, corrosion resistance, and biocompatibility, Sn is selected for incorporation into the Ti20Zr20Hf20Nb20Ta20 HEA. The effects of Sn alloying on the microstructures, mechanical properties, and corrosion behavior of the HEAs were investigated, and the corresponding mechanisms were also discussed.

2 Experimental

The alloy ingots with nominal compositions of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%, hereafter denoted as Sn3, Sn5, and Sn7, respectively) were prepared by arc-melting under a Ti-gettered purified argon atmosphere. The mixture of Sn and Ti was melted initially to prepare a pre-alloy and then re-melted with pure Zr, Nb, and Ta. The actual chemical compositions of the resulted alloys were examined by an energy-dispersive X-ray spectrometer (EDS) and are listed in Table 1. The microstructures were examined by a D/max2500PC X-ray diffraction (XRD, Cu Kα), an Apreo S LoVac scanning electron microscope (SEM), a JXA8100 electron probe X-ray microanalysis (EPMA) with an attached EDS, and a transmission electron microscope (TEM, Tecnai G2 F20). The TEM specimens were prepared by grounding the alloy slices to 50 μm in thickness and then ion-milling with a liquid-nitrogen specimen cooling stage.

Table 1 Chemical compositions of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%, denoted as Sn3, Sn5, and Sn7, respectively) alloys determined by EDS (at%)
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The compressive mechanical properties of these HEAs were studied by a material testing machine (SANS CMT5504) at a strain rate of 1 × 10–3 s−1 at room temperature. The specimens (2 mm × 2 mm × 4 mm) were cut from the middle part of the cross section of the alloy ingots by an electro-discharge machine, and the surfaces of these specimens were mechanically polished with the 2000-grit SiC sandpapers. The Young’s modulus (E) and nanoindentation hardness (H) of the alloys were calculated from the nanoindentation tests, using a nanoindenter (Keysight-Tech G200) with a standard Berkovich tip. The surfaces of the specimens (10 mm × 10 mm × 1 mm) for the nanoindentation were mechanically ground with the 2000-grit SiC sandpapers and then polished with a Buehler MasterMet suspension on a soft cloth. Indentation was performed with the peak load of 100 mN. Densities of these alloys were determined by the Archimedean principle using distilled water.

For investigating the corrosion behaviors of the HEA under the simulated physiological environment, electrochemical measurements were carried out in the Hank’s solution (prepared by dissolving 8.00 g NaCl, 0.40 g KCl, 0.14 g CaCl2, 0.20 g MgSO4·7H2O, 0.35 g NaHCO3, 0.06 g KH2PO4, 1.00 g C6H12O6, and 0.12 g Na2HPO4·12H2O in 1 L distilled water) at 310 K. Specimens (10 mm × 10 mm × 1 mm) were polished with 2000-grit SiC sandpapers and then cleaned in the ethanol and distilled water. The commercial Ti–6Al–4V alloy was used as a counterpart. The electrochemical experiments were performed in a three-electrode cell, using a saturated calomel reference electrode (SCE) and a platinum counter electrode with a Princeton Applied Research VersaSTAT III electrochemical workstation. Prior to the potentiodynamic polarization test, the corrosion specimen was immersed in the solution for 1800 s when the open-circuit potentials (OCP) achieved a steady state. The potentiodynamic polarization curves were recorded in a potential range from 50 mV below the OCPs to 1.6 V at a sweeping rate of 0.833 V·s−1. The corrosion rates of the alloys were calculated based on ASTM G102-89 with the equation:

$${\text{Corrosion rate }}({\text{mm}}\cdot{\text{year}}^{{ - {1}}} ) \, = { 3}.{27 } \times { 1}0^{{{-}{3}}} \times \left( {i_{{{\text{corr}}}} /\rho } \right) \times {\text{ EW}}$$
(1)

where icorr is the corrosion current density in μA·cm−2, ρ is the density in g·cm−3, and EW is the alloy equivalent weight, which is given by:

$${\text{EW}} = \, \left( {\sum \frac{{n_{i} f_{i} }}{{W_{i} }}} \right)^{{ - {1}}}$$
(2)

where ni is the valence of the ith element of the alloy, and fi is the mass fraction of the ith element in the alloy, and Wi is the atomic weight of the ith element in the alloy.

Mouse MC3T3-E1 pre-osteoblasts were used to estimate the in vitro biocompatibility of the HEAs. The specimens (Φ6 mm × 1 mm) were ground with 3000-grit SiC sandpapers and then washed in the ethanol and distilled water. Prior to the cell-culture experiments, all the specimens were sterilized by the exposure to an ultraviolet (UV) light for at least 3 h. The cell-culture experiments were performed according to the procedures described in Ref. [45]. The viability of cells after 24-h incubation was evaluated with a fluorescent dye (LIVE/DEAD Viability/Cytotoxicity Kit). The detailed procedures of staining cells were the same as those in a previous work [23]. The stained cells were viewed under a laser scanning confocal microscope (Olympus FV1200). At least triplicate specimens were used in all the cell-culture experiments for reproducibility. The results were statistically analyzed, using the Student’s t test.

3 Results and discussion

3.1 Microstructure

XRD patterns of the (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%) HEAs are presented in Fig. 1a. Only one group of diffraction peaks corresponding to the BCC phase is observed in the pattern of each alloy, indicating that these HEAs are composed of the single BCC solid solutions without intermetallic compounds. As shown in Fig. 1b, with the Sn contents increasing, the enlarged (110) peaks at 35°–39° tend to shift to the higher 2θ angle, indicating the decrease in the lattice constants of HEAs. The typical SEM backscattered-electron (BSE) images of the Sn3 and Sn7 HEAs, and the corresponding EDS mapping obtained by means of EPMA are displayed in Fig. 2. It is seen that the HEAs exhibit the typical dendritic microstructure, and the dendritic and the inter-dendritic regions are represented by bright and dark gray contrast, respectively. The addition of Sn leads to the changes in the morphology and dendritic size of the HEAs. With the amount of Sn increasing to x = 7, the microstructure of the HEA becomes coarse, as shown in Fig. 2b. From the corresponding EDS mapping in Fig. 2, it can be found that Nb and Ta are enriched in the bright dendritic regions (DR), while the dark inter-dendritic regions (ID) contain high concentrations of Ti, Zr, Hf, and Sn. The elemental segregation occurred due to the distribution of the elements during solidification. Ta and Nb with high melting temperatures of 3290 and 2750 K, respectively, prefer to form the dendritic arms, and Ti (1941 K), Zr (2128 K), Hf (2506 K), and Sn (505 K) with the relatively low melting temperatures are segregated into the ID, which is consistent with previous reports on the refractory HEAs. [17, 18, 26, 46]. The coarsening of the dendritic structure could be ascribed to the increasing content of Sn with the lowest melting temperature in this alloy system [24]. To further investigate the microstructures, TEM images and the corresponding selected area electron diffraction (SAED) patterns of the Sn3, Sn5, and Sn7 HEAs are displayed in Fig. 3. There are no precipitated phases in the Sn3 HEA, as shown in the bright-field TEM (BF-TEM) images (Fig. 3a). In the SAED patterns (the inset), only the diffraction spots of a BCC phase can be recognized, further confirming the formation of the single BCC solid solution with no other phases in the Sn3 HEA. With the increase in Sn content, nanoparticles distributed along the grain boundaries are found in the Sn5 and Sn7 HEAs, as shown in Fig. 3b, c. The two HEAs are mainly consisted of a BCC solid solution, which is in agreement with the XRD patterns. The nanoparticles are identified as a hexagonal-close-packed (HCP)-(Sn, Zr) ordered phase [42], which could not be detected in the XRD patterns due to the low volume fraction. With the Sn content increasing in the (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx HEAs, the size of the precipitates in the grain boundaries becomes larger.

Fig. 1
figure 1

a XRD patterns and b enlarged (110) diffraction peaks of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%) HEAs

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Fig. 2
figure 2

Typical SEM images and corresponding EDS mapping results of a Sn3 HEA and b Sn7 HEA

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Fig. 3
figure 3

TEM images and (inset) SAED patterns of a Sn3 HEA, b Sn5 HEA, and c Sn7 HEA

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Empirical parameters, such as mixing entropy (ΔSmix), mix enthalpy (ΔHmix), atomic-size difference (δ), and valence electron concentration (VEC), have been proposed to predict the phase formation in HEAs from the thermodynamic point of view [5, 7, 27]. The values of these parameters for the present HEAs are calculated and listed in Table 2. It is found that the values of the δ and VEC of these alloys are quite similar. With the addition of Sn, the values of ΔSmix increase, while the values of ΔHmix decrease. The negative values of ΔHmix for the Sn5 and Sn7 HEAs indicate that the chemical bonding between the components becomes stronger, and has a tendency to form compounds. Furthermore, Sn prefers to combine with Zr due to their strong interaction (ΔHSn-Zr = − 43 kJ·mol−1) [47], which could lead to the formation of the HCP-(Sn, Zr) ordered phase distributes in the grain boundaries.

Table 2 Mixing entropy (ΔSmix), mix enthalpy (ΔHmix), atomic-size difference (δ), and valence electron concentration (VEC) of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%) alloys
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3.2 Mechanical properties

Figure 4a shows the true compressive stress–strain curves of the (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at.%) HEAs. The values of the yield strengths (σ0.2) taken at a plastic strain of 0.2% and plastic strain (ε) are obtained and summarized in Table 3, in which the values of the Ti20Zr20Hf20Nb20Ta20 HEA are also shown for comparison [20,21,22]. It is notable that the strengths of the HEAs increased with the addition of Sn. The yield strengths of the Sn3 and Sn5 HEAs are gradually enhanced to reach 1068 and 1235 MPa, respectively, and the compressive strain slightly deceased to ~ 45% and ~ 42%. For the Sn7 HEA, the yield strength is further enhanced, but the plasticity dramatically reduced to ~ 30%. The improvement of yield strengths with the addition of Sn may be induced by solid solution strengthening and precipitation strengthening [8, 42, 48,49,50]. The plasticity of these HEAs exhibits a deteriorating tendency with the increase in the amount of Sn, which could be attributed to the weak interface strength between the precipitates and the main BCC phase [7, 51]. The surface of the Sn3 HEA after the compressive tests exhibited proper plastic flow and obvious buckling without visible microcracks based on the SEM observation. The Sn5 HEA also showed plastic flow, while some macrocracks were observed on the surface of the fractured samples along the direction of the main shear stress. The high yield strength and good plasticity of the novel Ti–Zr–Hf–Nb–Ta–Sn HEAs imply the potential of the alloys as biomaterials.

Fig. 4
figure 4

a True compressive stress–strain curves, b Representative load (P)—displacement (h) nanoindentation curves under nanoindentation tested with a Berkovich indenter, and c hardness (H) and reduced elastic modulus (E) of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%) HEAs

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Table 3 Density (ρ), yield strength (σ0.2), plastic strain (ε), hardness (H), Vickers hardness (Hv), and Young's modulus (E) of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 0, 3, 5, and 7 at%) HEAs
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The yield strengths of these HEAs theoretically estimated using the rule of mixture, i.e., (σ0.2)mix = Σci(σ0.2)i (where ci is the atomic fraction and (σ0.2)i is the yield strength of element i) for the Sn3, Sn5, and Sn7 HEAs are 219, 215, and 211 MPa, respectively, which are much lower than the values obtained from experiments. The solid solution strengthening effect in a single-phase HEA is generally considered on the basis of the misfit effects in the atomic size and modulus [8, 11, 18, 20]. In this case, HEA is taken as a pseudo-binary solid solution. The relative atomic size difference (δaij) and modulus difference (δGij) of the alloying element pairs as well as the calculated values of the lattice distortion (δai) and modulus distortion (δGi) for the Sn3 HEA are calculated and listed in Table 4. Owing to the large atomic size difference of Zr and Hf with other elements (δai ≈ 0.062), the HEA is considered as a pseudo-binary solid solution, defining the Ti, Nb, Ta, and Sn as the solvents and Zr plus Hf as the solutes. The strengthening contribution of atomic-size misfit is approximately quantified to be Δσa ≈ 252 MPa, using the equations in Ref. [20]. For calculating the contribution of modulus distortion, because the largest modulus misfit appeared in the vicinity of Sn (δGi ≈ − 0.832), the HEA is considered as another pseudo-binary solid solution with the Ti, Zr, Hf, Nb, and Ta as the solvents and Sn as the solute. The modulus distortion is roughly quantified to be ΔσG ≈ 531 MPa. Using the same calculating method, the strength contributions from lattice distortion (Δσa) and modulus distortion (ΔσG) and (σ0.2) mix for the Sn5 and Sn7 HEAs are also estimated and listed in Table 5, together with the experimental values. It can be found that the calculated yield strength (σC0.2) is close to the yield strength (σ0.2) obtained from experiments. This result suggests that the improvement of the yield strength of the these HEAs is attributed to the solid solution strengthening effect of the Sn addition. For the Sn5 and Sn7 HEAs, the contribution of precipitation strengthening to the yield strength could be negligible, since the precipitated HCP phases are in a small amount.

Table 4 Relative atomic size difference, δaij, and modulus difference, δGij (bold numbers), between atomic pairs in Sn3 HEA
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Table 5 Calculated strength contributions from (σ0.2)mix, estimated using rule of mixture, lattice distortion (Δσa), modulus distortion (ΔσG), calculated yield strength (σC0.2), and yield strength (σ0.2) obtained from experiments of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 0, 3, 5, and 7; at%) HEAs
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The Young’s modulus (E) and hardness (H) of the HEAs were tested by nanoindentation. The representative load (P)–displacement (h) nanoindentation curves are displayed in Fig. 4b, and the values of E and H obtained from Fig. 4b, using the Oliver-Pharr method [52], are presented in Fig. 4c. The Sn3 HEA exhibits a relatively low Young’s modulus of 80 GPa, similar to that of the equimolar Ti–Zr–Hf–Nb–Ta HEA [21]. The further addition of 5 at% and 7 at% Sn increases the E values of the HEAs (86–91 GPa), which could be due to the formation of HCP-(Sn, Zr) ordered phases in the grain boundaries [42]. The Young’s moduli of the present Ti–Zr–Hf–Nb–Ta–Sn HEAs are much lower than those of traditional metallic biomaterials, such as Co–Cr–Mo alloys (~ 240 GPa), 316L stainless steels (~ 210 GPa), and Ti–6Al–4V alloy (~ 110 GPa), and the low Young’s modulus is beneficial for alleviating the stress-shielding effect for the application as implant biomaterials. Moreover, for comparison, the H values of the HEAs are converted into the equivalent Vickers hardness (Hv) based on the relation of Hv = 0.9081 H for a Berkovich indentation tip based on ISO 14577-1:2016. As displayed in Table 3, the increase in the Sn content of the HEAs also improves the hardness of these HEAs, ranging from HV 315 to HV 390, which could be attributed to the solid solution hardening mechanism and the precipitation strengthening [18, 26, 42, 50]. Furthermore, the ratio of the hardness to Young’s modulus (H/E) of the Ti–Zr–Hf–Nb–Ta–Sn HEAs is in the range from 0.043 to 0.047, higher than those of the Ti20Zr20Hf20Nb20Ta20 HEA (~ 0.028) and the Ti–6Al–4V alloy (~ 0.026) [22], suggesting their good anti-wear property [53]. Overall, the combination of the low Young's moduli, high strengths, hardness and H/E values, and good plasticity of the Ti–Zr–Hf–Nb–Ta–Sn HEAs implies their potential for the biomedical application, such as dental and orthopedic implants. Among these alloys, the Sn3 HEA exhibits the lowest Young’s modulus (~ 80 GPa), high hardness (3.47 GPa), and compressive yield strength (~ 1068 MPa), together with good plasticity (~ 45%), implying the great potential as a biomaterial.

3.3 Corrosion behavior

To characterize the bio-corrosion behavior of the Ti–Zr–Hf–Nb–Ta–Sn HEAs, electrochemical measurements were taken at 310 K in Hank solution, and the Ti–6Al–4V alloys were used as the counterpart. Figure 5a presents the changes in OCPs with the immersion time of these alloys. During the immersion, the OCP values of the HEAs increase initially and achieve to be steady with the extending immersion time, implying the enhancement of the stability of the surface films in the solution. With the increase in the Sn content from 0 at% to 7 at%, the OCP after immersion for 1800 s of these HEAs slightly decreases from − 0.39 to − 0.43 V. It is suggested that the stability of the surface film is slightly weakened with the addition of Sn. The potentiodynamic–polarization curves of the HEAs and the Ti–6Al–4V alloy in the Hank’s solution are shown in Fig. 5b. All these HEAs are spontaneously passivated with low passive current densities (ipass) comparable to that of the Ti–6Al–4V alloy. For all these HEAs, no pitting occurs by the anodic polarization up to 1.6 V. Table 6 lists the corrosion parameters derived from the polarization curves. It is seen that with the Sn content increasing from 0 at% to 7 at%, the corrosion current density (icorr) and the corrosion potential (Ecorr) exhibit a slight increase and decrease tendency, respectively. The corrosion rates calculated from the icorr of these HEAs are on the same order of 1 × 10–4 mm·year−1, which is parallel to that of the Ti–6Al–4V alloy. These results indicate the good bio-corrosion resistance of the HEAs in the aggressive body fluid, which is favorable for the biomedical uses. The good corrosion resistance of the HEAs could be mainly attributed to the high concentration of Ta, Nb, Zr, Hf, and Ti, which could lead to the formation of protective passivation films on the alloys [22,23,24]. Sn oxides are also protective with the high potentials against chloride-ion attack, but they are less stable than the oxides of the valve metals (i.e., Ti, Nb, Ta, Hf, and Zr) [43, 44], which may lead to the increase in corrosion rate of the HEAs with high Sn amounts.

Fig. 5
figure 5

aChanges in open circuit potential with immersion time and b potentiodynamic-polarization curves of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%) HEAs and Ti–6Al–4V alloy in Hank’s solution at 310 K

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Table 6 Values of corrosion parameters derived from potentiodynamic–polarization curves of (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 0, 3, 5, and 7; at%) HEAs and Ti–6Al–4V alloy in Hank’s solution at 310 K
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3.4 In vitro biocompatibility

According to the aforementioned results, it is seen that the Sn3 HEA exhibited a good combination of the high yield strength and hardness, low Young’s modulus, and good bio-corrosion resistance. Thus, the in vitro biocompatibility of the Sn3 HEA was further evaluated through the direct cell-culture experiment, using the MC3T3-E1 cells compared with those of Ti20Zr20Hf20Nb20Ta20 HEA (denoted as Sn0) and the Ti–6Al–4V alloy. As indicated by the live/dead staining in Fig. 6a–c, numerous cells are adhered onto the surfaces of the HEAs and the Ti–6Al–4V alloy, and no dead/unhealthy cells are observed. It suggests that the alloys exhibited high viability and supported the initial adhesion of the MC3T3-E1 cells in a relatively short culture time of 24 h. Furthermore, the number of the cells attached onto the surfaces is presented in Fig. 6d. No statistical difference in the cell numbers on the HEAs and Ti–6Al–4V alloy is found, which indicates the good cell adhesion with high viability on the HEAs parallel to that on the Ti–6Al–4V alloy. The cell response to the Sn3 HEA is comparable with that to the Sn0 HEA, implying that the involvement of Sn in the present HEAs has no distinct effect on the biocompatibility of the HEA. Therefore, these results suggest that the (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)97Sn3 HEA possesses good in vitro biocompatibility, and the initial biosafety of the HEA is comparable to that of the Ti–6Al–4V alloy. From the perspective of biocompatibility, the surface films of these present HEAs are mainly composed of biocompatible elements Ti, Zr, Hf, Nb, and Ta [11, 19, 23, 24, 54]. Sn has also been used as a constituent element in Ti-based metallic biomaterials (e.g., Ti–Sn, Ti–Nb–Sn, Ti–Nb–Sn) with good biocompatibility that revealed by some in vitro cytotoxicity studies and animal implantation [40,41,42, 55], though pure Sn might affect the cell proliferation over time [56]. Furthermore, it has been reported that the biocompatibility of biomedical alloys is also influenced by their anti-corrosion properties [18, 23, 24]. The good corrosion resistance of the present HEAs in the corrosive body fluid is favorable for the biomedical application to inhibit the release of metallic ions into the tissue and guarantee the implant integrity [46, 54]. Thus, the present Ti–Zr–Hf–Nb–Ta–Sn HEAs exhibit good in vitro biocompatibility, implying their potential as biomaterials.

Fig. 6
figure 6

Live/dead staining of MC3T3-E1 cells on a Sn3 HEA, b Sn0 HEA, and c Ti–6Al–4V alloy, and d numbers of cells attached to HEAs and Ti–6Al–4V alloy after cell culture for 24 h (presented as percentages of number of cells attached to Ti–6Al–4V alloy)

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4 Conclusion

In the present study, (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)100-xSnx (x = 3, 5, and 7; at%) HEAs with good mechanical properties, corrosion resistance, and biocompatibility have been developed. The HEA with 3 at% Sn exhibits the single BCC structure. With the further increasing in Sn contents to x = 5 at% and 7 at%, the HEAs are composed of the BCC structure and the HCP-(Sn, Zr) ordered phase. These HEAs exhibit improved compressive yield strengths (1068–1259 MPa) and hardness (HV 315–HV 390) with the Sn contents increasing. Young's moduli of these HEAs are about 80–91 GPa, which are lower than that of the Ti–6Al–4V alloy. Among these HEAs, the (Ti0.2Zr0.2Hf0.2Nb0.2Ta0.2)97Sn3 HEA shows a good integration of mechanical properties, including a relatively low Young’s modulus of 80 GPa, high compressive yield strength up to ~ 1068 MPa, and a plastic strain of about 45%. Good bio-corrosion resistance of the HEAs is revealed by the passivation and no pitting in Hank’s solution. The HEA with 3 at% Sn can support the initial cell attachment with high viability, indicating the good biocompatibility. Combining the good mechanical properties, corrosion resistance, and biocompatibility, the Ti–Zr–Hf–Nb–Ta–Sn HEAs show great promise as potential biomaterials.