Effect of Niobium and Oxygen Contents on Microstructure and Mechanical Properties of α+β-Type Ti–(5–25)Nb–(0.5–1)O (mass%) Alloys for Biomedical Applications

Abstract

The microstructures, mechanical properties, and biocompatibilities of α+β-type Ti–Nb–O alloys with a wide range of Nb and oxygen contents were investigated. Ingots of Ti–(5–25)Nb–(0.5–1)O (mass%) alloys were prepared by arc melting and hot forged. The forged bars were heat-treated at 873–1373 K for 3.6 ks followed by quenching. The volume fraction of equiaxed α-phase (f_α) in the alloys was determined experimentally. The Ti–5Nb–yO and Ti–(15–20)Nb–yO alloys contained α′- and α′′-martensite, respectively, while the Ti–10Nb–yO alloys contained both α′- and α′′-martensite. Therefore, changing the Nb content alters the α′/α′′ ratio. The addition of oxygen increased the distribution coefficient of Nb and accelerated Nb distribution in the β-phase, which increased the boundary temperature for the formation of α′- and α′′-martensite in the Ti–5Nb–yO alloys. The Ti–5Nb–(0.5–0.75)O alloys with f_α value of 0.5 had higher elongation and strength and a lower elastic modulus than Ti–6Al–4V alloy. Ion elution from the Ti–5Nb–(0.5–0.75)O alloys into 0.1 M NaCl–0.1 M lactic acid solution was comparable to that of a Ti–6Al–4V extra low interstitial (ELI) alloy. The Ti–5Nb–(0.5–0.75)O alloys did not exhibit cytotoxicity, indicating their excellent biocompatibility. Thus, we propose that Ti–5Nb–(0.5–0.75)O alloys are suitable for use as low-cost α+β-type Ti alloys for biomedical applications.

1. Introduction

Titanium (Ti) and its alloys are used in aerospace components,^1,2) chemical plants,³⁾ and biomedical devices⁴⁾ because of their high specific strength,⁵⁾ excellent corrosion resistance,⁶⁾ and good biocompatibility.^7,8) α+β-type Ti alloys exhibit an excellent balance of strength and ductility. Their mechanical properties can be controlled by changing the phase fraction, shape, and grain size of the α-phase through thermomechanical processing.^9,10) α+β-type Ti alloys undergo a martensitic transformation upon quenching, wherein hexagonal close-packed (hcp) α′-martensite and orthorhombic α′′-martensite phases are formed from the β-phase.¹¹⁾ One of the factors that determine the formation of martensite during quenching is the concentration of β-stabilizing elements in the β-phase.¹²⁾ α′-martensite, which forms in alloys with lower concentrations of β-stabilizers, increases the hardness and tensile strength.¹³⁾ On the contrary, α′′-martensite, which requires a higher concentration of β-stabilizers than α′-martensite, decreases the hardness and elastic modulus.¹⁴⁾

Ti and its alloys have a high production cost.^15,16) Therefore, alloying with ubiquitous elements has been suggested as an effective method of reducing the cost.¹⁷⁾ In previous studies, oxygen has been used as an alloying element for producing α+β-type Ti alloys at a low cost, as oxygen is the main impurity in Ti sponge and has a high chemical affinity for Ti.^18–21) The Ti–1Fe–(0.3–0.35)O–(0.01–0.05)N (mass%) alloys have also been developed as low-cost α+β-type Ti alloys.^22–25) Iron (Fe) is another of the main impurity elements in Ti sponge. In this alloy, it acts as a β-stabilizing element. However, TiFe intermetallic compounds tend to form in Ti alloys that contain Fe when exposed to moderate temperatures for long periods.^24,25) The formation of these intermetallics decreases the tensile strength and elongation of the alloy.

In our previous work, we explored the use of vanadium (V) as an alloying element and investigated the microstructure and mechanical properties of α+β-type Ti–V–O alloys.^12,26) Compared with V, niobium (Nb) has low cytotoxicity and is already employed as an alloying element in certain biomedical β-type Ti alloys.²⁷⁾ Furthermore, Nb is a β-isomorphous-type element and thus does not form intermetallic compounds with Ti.²⁸⁾ Because β-type Ti–Nb binary alloys show high superelasticity and a strong shape memory effect,^29–31) Ti–Nb–O alloys are expected to be suitable for use as biomedical Ti alloys. However, to the best of our knowledge, the microstructure and mechanical properties of α+β-type Ti–Nb–O alloys with a wide range of Nb and oxygen contents for biomedical applications have not yet been reported.

The objective of this study was to investigate the effects of Nb and oxygen on the microstructure and mechanical properties of α+β-type Ti–Nb–O alloys to develop low-cost α+β-type Ti alloys. To evaluate the suitability of the alloys for use in biomedical applications, their cytotoxicity and ion elution properties in a 0.1 M NaCl–0.1 M lactic acid solution were evaluated. In this study, the Nb and oxygen contents were varied up to 25 and 1 mass%, respectively. Hereafter, all chemical compositions are given in mass%.

2. Experimental Procedures

2.1 Fabrication and heat treatment of specimens

Commercially pure (CP) Ti plates (JIS Grade 2, UEX Ltd., Tokyo, Japan, O: 0.1038% (1038 mass ppm)) and TiO₂ powder (99.5%, Kanto Chemical Co. Inc., Tokyo, Japan) were melted using a non-consumable-electrode argon (Ar) arc melting furnace to fabricate a master Ti–2.4%O binary alloy. Next, the master alloy, CP Ti plates, and Nb plates (99.9%, Nilaco Corp., Tokyo, Japan) were arc-melted to fabricate Ti–(5–25)%Nb–(0.5–1)%O alloy ingots. The oxygen content in the alloys was measured by inert gas fusion-infrared absorption (ONH836, LECO, MI), and the Nb content was measured using inductively coupled plasma-mass spectrometry (ICP-MS, Agilent8800, Agilent Technologies, CA). The chemical compositions of the alloys are shown in Table 1. Hereafter, the alloys are notated as Ti–xNb–yO using the alloy composition in mass%. After melting, the ingots were forged in the β-region at 1373 K to form bars with a diameter of 16 mm, followed by air cooling to room temperature. Then, the alloys were forged in the α+β region at 1023 K to form bars with a diameter of 8 mm, followed by air cooling to room temperature. During the forging process, the alloys were re-heated to prevent a decrease in their temperature. The forged alloy bars were cut into sections (diameter of 8 mm and length of 40 mm), and the oxide scale on the specimen surface was removed by machining. After forging, heat treatments were performed under Ar flow at 873–1373 K for 3.6 ks, followed by quenching in ice water.

Table 1 Chemical composition of Ti–xNb–yO alloys used in this study (mass%).

2.2 Specimen analysis

The constituent phases of the heat-treated alloys were identified by X-ray diffraction (XRD; D8 ADVANCE, Bruker AXS K.K., Karlsruhe, Germany). Their microstructures were observed by scanning electron microscopy (SEM) in backscattered electron (BSE) mode (XL30FEG, Royal Philips Electronics Inc., Amsterdam, Netherlands). The volume fraction of the equiaxed α-phase (f_α) was calculated from five SEM images (×2000 magnification) using image analysis software (ImageJ, NIH, US). The value of f_α was calculated as the area fraction of the equiaxed α-phase observed in the SEM images according to eq. (1).^12,26)   

\begin{equation} f_{\alpha} = \frac{S_{\alpha}}{S_{\text{Total}}} \end{equation} (1)

where S_α is the area of the equiaxed α-phase in the SEM image, and S_Total is the total image area.

Further microstructural observations were conducted using transmission electron microscopy (TEM; JEM-ARM200F, JEOL Ltd., Tokyo, Japan). The Nb contents of the α- and β-phases were determined by field-emission electron probe microanalysis (FE-EPMA; JXA-8530F, JEOL Ltd., Tokyo, Japan) using CP Ti and Ti–(5, 10, 15, 20)Nb–(0.5, 0.7, 1)O alloys as reference materials. The reference materials were heat-treated above the β-transus temperature (T_β) to obtain a single phase. Measurements were performed at 10 positions for each phase, and the average values and standard deviations were calculated.

The Vickers hardness of the heat-treated alloys was determined using a Vickers microhardness tester (HM-102, Mitutoyo Co., Kanagawa, Japan) under a load of 50 gf. Uniaxial tensile tests were conducted at room temperature using a universal material testing system (RTF-1325, A&D Co. Ltd., Tokyo, Japan). The tensile specimens were machined to a gauge diameter and length of 3 and 10 mm, respectively.^32,33) The strain rate was set to 5.0 × 10⁻⁴ s⁻¹,^12,26) and the tensile strength, 0.2% proof stress, and total elongation were measured. The fracture surfaces were imaged by SEM to measure the reduction in area. The elastic moduli of the specimens were measured using the free resonance vibration method (JE-RT3, Nihon Techno-Plus Co. Ltd., Osaka, Japan). The tests were performed on three specimens for each heat-treatment condition, and the average values and standard deviations were calculated.

2.3 Evaluation of corrosion resistance

To evaluate the corrosion resistance of the alloys, they were subjected to static immersion tests in 0.1 M NaCl–0.1 M lactic acid solution (pH: 2.3 ± 0.1) in accordance with ISO 10271.³⁴⁾ The as-forged alloys were cut into coin-shaped specimens, each with a diameter of 7 mm and thickness of 2 mm (surface area (S): 121 mm²), and the specimen surfaces were polished using abrasive papers with grit numbers of up to #1500. The polished specimens were washed ultrasonically in ethanol and ultrapure water for 300 s each before the static immersion tests. The immersion tests were performed in triplicate. Each test specimen was placed in a Teflon vessel with 10 mL of the 0.1 M NaCl–0.1 M lactic acid solution. The vessel was then placed in a shaking water bath at 310 K for 604.8 ks (1 week) and shaken at a speed of 60 rpm. The concentrations of Ti and Nb ions eluted from the alloy specimen were measured by ICP-MS. The Ti–6Al–4V extra low interstitial (ELI) alloy was also evaluated for comparing the fabricated alloy specimens with a commercially available biomedical α+β-type Ti alloy.

2.4 Evaluation of cytotoxicity

Direct and indirect cytotoxicity tests were performed according to ISO 10993-5.³⁵⁾ The detailed procedure for the tests is described elsewhere.³⁶⁾ Fibroblasts from the lungs of Chinese hamsters (V79 cells) were used for the evaluations. Coin-shaped as-forged specimens with diameters of 15 mm and thicknesses of 2 mm were used as specimens. They were washed ultrasonically in acetone and ultrapure water for 300 s each and autoclaved at 394 K for 1.26 ks (21 min).

For the direct cytotoxicity tests, five specimens corresponding to each alloy composition were placed separately in a 24-well plate. Next, V79 cells (50 cells·mL⁻¹) in 1 mL of minimum essential medium (MEM) 10 (Eagle’s MEM containing 10% fetal bovine serum and 100 nM sodium pyruvate) were seeded on each specimen. Wells without alloy specimens were used as controls. The cells in the 24-well plates were then incubated for 7 days at 310 K in 5% CO₂.

For the indirect cytotoxicity tests, five specimens corresponding to each alloy composition were immersed in 20 mL of MEM 10 without any V79 cells and incubated for 24 h at 310 K in 5%CO₂. After incubation, 1 mL of the extract solution was added to each well of five separate 24-well plates. The V79 cells were seeded in all the plates at a rate of 50 cells·mL⁻¹, and wells without extract solution were used as controls. The cells were incubated in the extracts for 7 days at 310 K in 5%CO₂.

After the completion of the incubation process, the viable cell colonies were stained with 5% Giemsa solution. For the direct cytotoxicity tests, the cell colonies on the coin-shaped specimens and those in the control well plates were counted. For the indirect cytotoxicity tests, the cell colonies in the wells were counted.

The results of the cytotoxicity measurements were statistically analyzed using one-way ANOVA and Tukey’s multiple comparison tests.

3. Results and Discussion

3.1 Microstructure

The SEM/BSE images of the heat-treated Ti–5Nb–0.5O alloys are shown in Fig. 1. Equiaxed α-phase grains were observed after heat-treatment between 973 and 1173 K, while no equiaxed grains were detected after heat-treatment at 1223 K. The volume fraction of equiaxed α-phase (f_α) in each alloy was calculated from the SEM images. The f_α values are plotted in Fig. 2 as a function of the heat-treatment temperature (approach curve). In the cases where no equiaxed α-phase was observed in the alloy, such as Fig. 1(f), the value of f_α was taken to be zero. Here, the approach curves were obtained as smooth curves using all values of f_α except f_α = 0 for each alloy. The value of f_α decreased with an increase in the heat-treatment temperature. The β-transus temperatures (T_β) of the alloys were estimated by extrapolating the approach curves to f_α = 0. Figure 3 depicts the relationship between the estimated T_β values (Fig. 2) and the Nb content of the alloys. The T_β values decreased with an increase in the Nb content and a decrease in the oxygen content. The T_β value of the Ti–Nb–O alloy system was obtained using the multiple regression analysis of these results and determined as   

\begin{equation} T_{\beta}\ (\text{K}) = 1155 + 234[\text{O}] - 8.1[\text{Nb}] \end{equation} (2)

where [O] and [Nb] are the concentrations of oxygen and Nb (mass%) in the alloys, respectively. The coefficient of determination, R², was 0.988, and P values of oxygen and Nb were <0.001. The results obtained using this equation were consistent with the measured values within ±20 K.

Fig. 1

Microstructure of Ti–5Nb–0.5O alloy after heat-treatment at (a) 973, (b) 1023, (c) 1073, (d) 1123, (e) 1173, and (f) 1223 K.

Fig. 2

Volume fractions of equiaxed α-phase of Ti–xNb–yO alloys as a function of heat-treatment temperature. (a) y = 0.5, (b) y = 0.75, and (c) y = 1.

Fig. 3

β-transus temperatures (T_β) of Ti–xNb–yO alloys as determined from the approach curves.

The XRD patterns of the heat-treated Ti–(5–20)Nb–0.5O alloys are shown in Fig. 4. The dotted lines in the figure indicate the position of T_β for each alloy. For the Ti–5Nb–0.5O alloy, peaks related to the hcp phase were detected at all heat-treatment temperatures, while a peak related to the body-centered cubic (bcc) phase was detected after heat-treatment at 973 and 1023 K. Although the heat-treatment temperature of 1223 K is higher than the T_β value of this alloy, only hcp-phase-related peaks were detected in the XRD pattern of the alloy. An acicular transformed structure was observed in the alloy heat-treated at 1223 K (Fig. 1(f)); therefore, it is likely that the β-phase existed at 1223 K and transformed to hcp α′-martensite during quenching. Since the SEM images of the alloys heat-treated at between 1073 and 1173 K show both the equiaxed α-phase and acicular transformed structure, the hcp-phase-related peaks in the XRD patterns were considered to be α-phase and/or α′-martensite. Because of the abovementioned reasons, the peaks related to bcc and hcp structures were probably ascribable to the β-phase and α-phase and/or α′-martensite, respectively. In the case of the Ti–(10–20)Nb–0.5O alloys, the peaks corresponding to α′′-martensite were detected at higher heat-treatment temperatures.

Fig. 4

XRD patterns of heat-treated Ti–xNb–0.5O alloys. (a) x = 5, (b) x = 10, (c) x = 15, and (d) x = 20.

The Nb content and temperature region for the formation of α′- and α′′-martensite, as determined from the XRD analyses, are shown in Fig. 5. The conditions under which martensite was not formed are marked as “×”. The dotted and solid lines in the figure are the T_β and boundary lines for the formation of α′- and α′′-martensite, respectively. In this study, α′- and α′′-martensite structures were observed in the Ti–5Nb and Ti–(15–20)Nb alloys, respectively, irrespective of the oxygen content. In the Ti–10Nb–yO alloys, a temperature region for the formation of α′+α′′-martensite existed, and the boundary temperature increased with increasing oxygen content.

Fig. 5

Nb content and heat-treatment temperature for formation of α′- and α′′-martensite in Ti–xNb–yO alloys. (a) y = 0.5, (b) y = 0.75, and (c) y = 1.

For Ti–Nb binary alloys, α′′-martensite forms during quenching when the Nb content in the β-phase is greater than 12.7³⁷⁾–13.1³⁸⁾%, while α′-martensite forms if the Nb content is lower than this. The Nb contents of the equiaxed α- and β-phases of the Ti–10Nb–(0.5–1)O alloys, as measured by FE-EPMA, are shown in Fig. 6 as a function of the heat-treatment temperature. The dotted lines in the figure represent the Nb content of the alloys, as measured by ICP-MS. The solid bands represent the reported minimum Nb content for the formation of α′′-martensite (C_Nb = 12.7³⁷⁾–13.1³⁸⁾%). The Nb content in the equiaxed β-phase increased with a decrease in the heat-treatment temperature, while that in the equiaxed α-phase remained almost constant. For the Ti–10Nb–0.5O alloy (Fig. 6(a)), the Nb content of the β-phase increased from 12.2% to 14.2% as the heat-treatment temperature decreased from 1123 to 1073 K. Because the Nb content of the equiaxed β-phase was below 12.7%, α′-martensite formed in the Ti–10Nb–0.5O alloy after heat-treatment at 1123 K. In contrast, α′′-martensite formed after heat-treatment at 1073 K because the Nb content in the equiaxed β-phase (14.2%) was sufficient for α′′-martensite to form. These results are similar to those for the alloys with other oxygen contents. The distribution of Nb in the equiaxed β- and α-phase could explain the formation of α′- and α′′-martensite in this alloy system (Fig. 5). The oxygen content in each grain was measured using FE-EPMA. The measured oxygen contents in the β-phase were approximately 0.5, 0.5, and 0.7 mass% for Ti–10Nb–0.5O, 0.75O, and 1O alloys, respectively, where α′′-martensite was detected. The heat-treatment conditions with α′′-martensite formation (Fig. 5) were in good agreement with the Nb content in the β-phase, which was higher than 13 mass% (Fig. 6) even with such a significant difference in the oxygen content from 0.5 to 0.7 mass% in the β-phase. Wang et al.³⁹⁾ reported that oxygen did not change the phase constitution of Ti–15 mass% Nb alloy, while it suppressed the α′′-martensite and ω phases in the case of Ti alloys with higher Nb content of 32 mass% (higher β-phase stability). These results suggest that α′′-martensite formation was mainly dependent on the Nb content in the β-phase for the Ti–10Nb alloys and independent of the oxygen content.

Fig. 6

Nb content in equiaxed α- and β-phase grains of Ti–10Nb–yO alloys as a function of heat-treatment temperature. (a) y = 0.5, (b) y = 0.75, and (c) y = 1.

Next, we discuss the reason why the α′/α′′ phase boundary temperature increased with increasing oxygen content in the Ti–10Nb–yO alloys. The Nb distribution coefficients (k_Nb) were calculated from the Nb contents of each phase using the results given in Fig. 6, and are shown in Fig. 7 as a function of the volume fraction of equiaxed α-phase (f_α). The value of k_Nb is defined by eq. (3).   

\begin{equation} k_{\text{Nb}} = C_{\text{Nb},\alpha}/C_{\text{Nb},\beta} \end{equation} (3)

where C_Nb,α and C_Nb,β represent the Nb contents of the equiaxed α- and β-phases, respectively. At the same f_α value (e.g., 0.4), the k_Nb value increased with an increase in the oxygen content. This indicates that oxygen accelerates the distribution of Nb to the β-phase, resulting in an increase in the α′/α′′ phase boundary temperature.

Fig. 7

Distribution coefficient of Nb of Ti–10Nb–yO alloys as a function of the volume fraction of α-phase (f_α).

3.2 Mechanical properties

Figure 8 shows the relationship between the heat-treatment temperature and the Vickers hardness of the alloys. For the Ti–10Nb–0.5O and Ti–15Nb–0.5O alloys, the Vickers hardness increased sharply when heat-treated at 873–973 K and decreased as the heat-treatment temperature increased above 1023 K (Fig. 8(a)). For the other compositions, the Vickers hardness slightly increased with increasing heat-treatment temperature. TEM dark-field images and electron diffraction patterns of the 923 and 973 K-heat-treated Ti–10Nb–0.5O alloys are shown in Fig. 9. Fine precipitates were present in the β-phase of both specimens, as evident from the dark-field images (Figs. 9(a) and (c)). From the electron diffraction patterns of the β-phase (Figs. 9(b) and (d)), it was confirmed that these precipitates were of an athermal ω-phase (ω_ath), which is likely to have formed during quenching. Because ω_ath is known to increase the hardness of Ti alloys,^14,29) the marked increase in the hardness of the 923 and 973 K-heat-treated Ti–10Nb–0.5O alloys was due to the formation of ω_ath precipitates. The formation of ω_ath was also observed with the XRD analysis under these conditions. Ti alloys are typically used under conditions where ω_ath is not formed. In Ti–xNb binary alloys, ω_ath is detected in the alloys with Nb contents of more than 10 mass%, and the amount of ω_ath increases with increasing Nb content.⁴⁰⁾ Paton and Williams⁴¹⁾ reported that oxygen, which is an interstitial element in Ti alloys, suppresses the formation of ω_ath by limiting the displacement of atoms. In this study, the formation of ω_ath was not observed in the alloys with higher oxygen content, such as 0.75 and 1 mass%, with the XRD analysis (shown in Fig. A1).

Fig. 8

Vickers hardness of Ti–xNb–yO alloys as a function of heat-treatment temperature. (a) y = 0.5, (b) y = 0.75, and (c) y = 1.

Fig. 9

(a), (c) TEM dark-field images and (b), (d) electron diffraction patterns of Ti–10Nb–0.5O alloys heat-treated at (a), (b) 923 and (c), (d) 973 K.

The mechanical properties of α+β-type Ti alloys can be altered by varying the f_α value. Therefore, α+β processing is commonly adopted. Generally, the α+β processing temperature for α+β-type Ti alloys is ≤40 K below β-transus temperature (T_β),⁴²⁾ where the f_α value is approximately 0.5. In light of these considerations, we evaluated the tensile properties of the alloys with f_α values of approximately 0.5. The heat-treatment temperatures and f_α values used to evaluate the alloys are listed in Table 2. Under these conditions, no ω_ath was observed.

Table 2 Heat-treatment temperatures (T_H) used for evaluating mechanical properties of Ti–xNb–yO alloys and volume fractions of equiaxed α-phase (f_α).

The results are shown in Fig. 10. The alloys had a tensile strength of more than 950 MPa, regardless of the Nb and oxygen contents. Furthermore, the tensile strength and 0.2% proof stress increased with the oxygen content, while the total elongation decreased owing to oxygen solid-solution strengthening.⁴³⁾ Similarly, the total elongation of the Ti–xNb–0.75O alloys decreased with increasing Nb content.

Fig. 10

Tensile strength, 0.2% proof stress, total elongation, and reduction in area of heat-treated Ti–xNb–yO alloys.

The tensile properties of the α+β-type Ti–Nb–O alloys were compared with those of Ti–6Al–4V alloys.^6,44–50) The tensile strength and total elongation of these alloys are shown in Fig. 11. The Ti–5Nb–0.5O and Ti–5Nb–0.75O alloys exhibited superior elongation to the Ti–6Al–4V alloy. These results indicate that the Ti–5Nb–(0.5–0.75)O alloys are strong candidates for use as α+β-type Ti alloys with a high oxygen content, given their excellent balance of strength and elongation.

Fig. 11

Comparison of tensile strength and total elongation of Ti–xNb–yO alloys and Ti–6Al–4V alloy.^6,44–50)

The elastic moduli of the Ti–5Nb–yO and Ti–xNb–0.75O alloys are shown in Fig. 12. For the alloys with the 5% Nb, the elastic modulus increased slightly from 94 to 97 GPa as the oxygen content increased from 0.5% to 1%. The elastic moduli of the Ti–(10–20)Nb–0.75O alloys were approximately 90 GPa. For an alloy to be suitable for use in bone-substitute devices such as the stem of artificial hip joints and bone plates, it must exhibit a similar elastic modulus to that of bone to help prevent stress shielding of the bone.⁵¹⁾ From this perspective, the Ti–5Nb–0.5O and Ti–5Nb–0.75O alloys are suitable because they have a lower elastic modulus than that of Ti–6Al–4V ELI alloy (110 GPa).

Fig. 12

Elastic moduli of heat-treated Ti–xNb–yO alloys.

3.3 Biocompatibility of the Ti–5Nb–(0.5–0.75)O alloys

Owing to their microstructure, mechanical properties, and elastic modulus, the Ti–5Nb–(0.5–0.75)O alloys are good candidates for low-cost α+β-type Ti alloys. In order to evaluate the feasibility of using the fabricated alloys in biomedical applications, their corrosion resistances in simulated body fluid were evaluated. Figure 13 shows the amounts of metallic ions eluted per unit area from the Ti–5Nb–0.5O and Ti–5Nb–0.75O alloys into a 0.1 M NaCl–0.1 M lactic acid solution. The ion-elution values for a commercial biomedical Ti–6Al–4V ELI alloy are also shown for comparison. The amount of Ti eluted from the Ti–5Nb–0.5O and Ti–5Nb–0.75O alloys was comparable to that from the Ti–6Al–4V ELI alloy. Moreover, the amount of Nb eluted from the Ti–5Nb–0.5O and Ti–5Nb–0.75O alloys was approximately the same as the amount of Al and V eluted from the Ti–6Al–4V ELI alloy. The results of the immersion tests indicate that the Ti–5Nb–0.5O and Ti–5Nb–0.75O alloys have comparable corrosion resistance in simulated body fluid to Ti–6Al–4V ELI alloy.

Fig. 13

Mass of ions eluted per unit area from Ti–5Nb–(0.5–0.75)O and Ti–6Al–4V ELI alloys into 0.1 M NaCl–0.1 M lactic acid solution over 1 week.

The numbers of viable V79 cells on the Ti–5Nb–0.5O and Ti–6Al–4V ELI alloys, as measured by direct and indirect cytotoxicity tests, are shown in Fig. 14. For both Ti–5Nb–0.5O and Ti–6Al–4V ELI alloys, there were no significant differences in the number of viable cells as compared with that for the control (blank) tests, indicating that these alloys are not cytotoxic. Because the half-maximal inhibitory concentration (IC50) of Nb ions is 1000-times smaller than that of V ions,⁵²⁾ the Ti–5Nb–0.5O alloy should show a lower or similar level of cytotoxicity compared to that of the Ti–6Al–4V ELI alloy.

Fig. 14

Number of viable V79 cells as measured by direct and indirect cytotoxicity tests. n.s.: not significant.

The corrosion resistance and cytotoxicity tests confirmed that the Ti–5Nb–(0.5–0.75)O alloys are suitable for use as low-cost α+β-type Ti alloys for biomedical applications. In future studies, we plan to investigate how the oxygen content and heat-treatment of Ti–5Nb alloys affect their microstructure, mechanical properties, and cold workability, to develop a precise alloy design strategy.

4. Conclusions

The effects of the Nb and oxygen contents of α+β-type Ti–Nb–O alloys on their microstructure, mechanical and physical properties, and biocompatibility were investigated. The obtained results are summarized as follows:

1. (1)    The formation of α′-martensite was observed in the Ti–5Nb–yO alloys after quenching. An increase in the Nb content to 10 mass% led to the formation of α′- and α′′-martensite. Further increases in the Nb content to 15 and 20 mass% resulted in the formation of α′′-martensite. The addition of oxygen increased the distribution coefficient of Nb and accelerated Nb distribution in the β-phase, which in turn led to an increase in the boundary temperature for the formation of α′- and α′′-martensite in the Ti–10Nb–yO alloys.
2. (2)    The Ti–5Nb–(0.5–0.75)O alloys exhibited greater elongation than commercial Ti–6Al–4V alloy. The Ti–5Nb–(0.5–0.75)O alloys had an elastic modulus of 95 GPa, which is lower than that of the Ti–6Al–4V ELI alloy.
3. (3)    Ion elution from the Ti–5Nb–(0.5–0.75)O alloys into 0.1 M NaCl–0.1 M lactic acid solution was comparable to that from the Ti–6Al–4V ELI alloy. Furthermore, the Ti–5Nb–0.5O alloy did not exhibit cytotoxicity, indicating that it has excellent biocompatibility.
4. (4)    Because the Ti–5Nb–(0.5–0.75)O alloys exhibit an excellent balance of strength and ductility, as well as a lower elastic modulus and comparable biocompatibility to that of Ti–6Al–4V ELI alloy, we propose that they are ideal for use as low-cost α+β-type Ti alloys for biomedical applications.

Acknowledgments

The authors would like to thank Dr. K. Kobayashi of Tohoku University for help with the TEM analyses. This study was partially supported by Grants-in-Aid for Science Research grant numbers 20H02448 and 21H04603 from the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan, and The Light Metal Educational Foundation, Inc.

REFERENCES

• 1)   R.R.  Boyer: Mater. Sci. Eng. A 213 (1996) 103–114. doi:10.1016/0921-5093(96)10233-1
• 2)   Z.  Liu,  B.  He,  T.  Lyu and  Y.  Zou: JOM 73 (2021) 1804–1818. doi:10.1007/s11837-021-04670-6
• 3)   I.V.  Gorynin: Mater. Sci. Eng. A 263 (1999) 112–116. doi:10.1016/S0921-5093(98)01180-0
• 4)   M.  Kaur and  K.  Singh: Mater. Sci. Eng. C 102 (2019) 844–862. doi:10.1016/j.msec.2019.04.064
• 5)   M.  Peters,  J.  Kumpfert,  C.H.  Ward and  C.  Leyens: Adv. Eng. Mater. 5 (2003) 419–427. doi:10.1002/adem.200310095
• 6)   K.  Maeda,  S.  Suzuki,  K.  Ueda,  T.  Kitashima,  S.Kr.  Bhattacharya,  R.  Sahara and  T.  Narushima: J. Alloy. Compd. 776 (2019) 519–528. doi:10.1016/j.jallcom.2018.10.291
• 7)   M.  Niinomi,  T.  Hanawa and  T.  Narushima: JOM 57 (2005) 18–24. doi:10.1007/s11837-005-0076-3
• 8)   J.  Wu,  K.  Ueda and  T.  Narushima: Mater. Sci. Eng. C 109 (2020) 110599. doi:10.1016/j.msec.2019.110599
• 9)   C.  Ouchi: Bull. Japan Inst. Metals 25 (1986) 672–679. doi:10.2320/materia1962.25.672
• 10)   S.L.  Semiatin,  V.  Seetharaman and  I.  Weiss: JOM 49 (1997) 33–39. doi:10.1007/BF02914711
• 11)   Y.  Murakami: Tetsu-to-Hagané 73 (1987) 420–426. doi:10.2355/tetsutohagane1955.73.3_420
• 12)   K.  Ueda,  T.  Kobayashi and  T.  Narushima: J. Japan Inst. Met. Mater. 80 (2015) 60–65. doi:10.2320/jinstmet.JB201503
• 13)   H.  Matsumoto,  H.  Yoneda,  K.  Sato,  S.  Kurosu,  E.  Maire,  D.  Fabregue,  T.J.  Konno and  A.  Chiba: Mater. Sci. Eng. A 528 (2011) 1512–1520. doi:10.1016/j.msea.2010.10.070
• 14)   C.M.  Lee,  C.P.  Ju and  J.H.  Chern Lin: J. Oral Rehabil. 29 (2002) 314–322. doi:10.1046/j.1365-2842.2002.00825.x
• 15)   F.H.  Froes: JOM 50 (1998) 15. doi:10.1007/s11837-998-0407-2
• 16)   F.H.  Froes,  H.  Friedrich,  J.  Kiese and  D.  Bergoint: JOM 56 (2004) 40–44. doi:10.1007/s11837-004-0144-0
• 17)   M.  Niinomi: J. Japan Inst. Metals 75 (2011) 21–28. doi:10.2320/jinstmet.75.21
• 18)   O.  Takeda,  T.  Ouchi and  T.H.  Okabe: Metall. Mater. Trans. B 51 (2020) 1315–1328. doi:10.1007/s11663-020-01898-6
• 19)   O.  Takeda and  T.H.  Okabe: JOM 71 (2019) 1981–1990. doi:10.1007/s11837-018-3278-1
• 20)   M.  Watanabe,  F.  Sato,  K.  Ueda,  D.  Matsuwaka and  T.  Narushima: Mater. Trans. 58 (2017) 613–618. doi:10.2320/matertrans.M2017003
• 21)   M.  Watanabe,  F.  Sato,  R.  Abe,  K.  Ueda,  D.  Matsuwaka,  F.  Kudo and  T.  Narushima: Metall. Mater. Trans. B 50 (2019) 1553–1558. doi:10.1007/s11663-019-01608-x
• 22)  H. Fujii, S. Soeda, M. Hanaki and H. Okano: Titanium ‘95 Science and Technology: Proc. 8th World Conf. on Titanium, ed. by P.A. Blenkinsop, W.J. Evans and H.M. Flower, (UK, TIM, 1996) pp. 2309–2316.
• 23)   M.  Koike,  C.  Ohkubo,  H.  Sato,  H.  Fujii and  T.  Okabe: Mater. Sci. Eng. C 25 (2005) 349–356. doi:10.1016/j.msec.2005.04.002
• 24)  H. Fujii and T. Maeda: Nippon Steel & Sumitomo Metal Technical Report 106 (2014) 16–21.
• 25)   H.  Fujii: Titanium Japan 51 (2002) 33–37.
• 26)   M.  Omiya,  K.  Ueda and  T.  Narushima: Mater. Trans. 58 (2017) 1250–1256. doi:10.2320/matertrans.L-M2017825
• 27)   D.  Kuroda,  M.  Niinomi,  M.  Morinaga,  Y.  Kato and  T.  Yashiro: Mater. Sci. Eng. A 243 (1998) 244–249. doi:10.1016/S0921-5093(97)00808-3
• 28)   M.  Ikeda and  M.  Ueda: J. JILM 67 (2017) 538–544.
• 29)   H.Y.  Kim,  Y.  Ikehara,  J.I.  Kim,  H.  Hosoda and  S.  Miyazaki: Acta Mater. 54 (2006) 2419–2429. doi:10.1016/j.actamat.2006.01.019
• 30)   S.  Miyazaki,  H.Y.  Kim and  H.  Hosoda: Mater. Sci. Eng. A 438–440 (2006) 18–24. doi:10.1016/j.msea.2006.02.054
• 31)   H.Y.  Kim,  S.  Hashimoto,  J.I.  Kim,  H.  Hosoda and  S.  Miyazaki: Mater. Trans. 45 (2004) 2443–2448. doi:10.2320/matertrans.45.2443
• 32)   K.  Ueki,  K.  Ueda and  T.  Narushima: Metall. Mater. Trans. A 47 (2016) 2773–2782. doi:10.1007/s11661-016-3488-5
• 33)   K.  Ueki,  K.  Ueda,  M.  Nakai,  T.  Nakano and  T.  Narushima: Metall. Mater. Trans. A 49 (2018) 2393–2404. doi:10.1007/s11661-018-4597-0
• 34)  ISO 10271: 2020.
• 35)  ISO 10993-5: 2009.
• 36)   O.  Gokcekaya,  K.  Ueda,  K.  Ogasawara,  H.  Kanetaka and  T.  Narushima: Mater. Sci. Eng. C 75 (2017) 926–933. doi:10.1016/j.msec.2017.02.059
• 37)   H.  Kim,  H.  Hosoda and  S.  Miyazaki: J. JILM 55 (2005) 613–617. doi:10.2464/jilm.55.613
• 38)   D.L.  Moffat and  D.C.  Larbalestier: Metall. Mater. Trans. A 19 (1988) 1677–1686. doi:10.1007/BF02645135
• 39)   J.  Wang,  W.  Xiao,  L.  Ren,  Y.  Fu and  C.  Ma: Mater. Charact. 176 (2021) 111122. doi:10.1016/j.matchar.2021.111122
• 40)   M.-K.  Han,  J.-Y.  Kim,  M.-J.  Hwang,  H.-J.  Song and  Y.-J.  Park: Materials 8 (2015) 5986–6003. doi:10.3390/ma8095287
• 41)   N.E.  Paton and  J.C.  Williams: Scr. Metall. 7 (1973) 647–649. doi:10.1016/0036-9748(73)90229-9
• 42)   T.  Furuhara: J. JILM 68 (2018) 40–47.
• 43)   Z.  Liu and  G.  Welsch: Metall. Trans. A 19 (1988) 527–542. doi:10.1007/BF02649267
• 44)   H.  Arabi and  M.  Ketabchi: Mater. Sci. Technol. 35 (2019) 1735–1741. doi:10.1080/02670836.2019.1645389
• 45)   P.  Ahmadian and  M.  Morakabati: Trans. Indian Inst. Met. 73 (2020) 1301–1309. doi:10.1007/s12666-020-01957-1
• 46)   M.K.  Dunstan,  J.D.  Paramore,  Z.Z.  Fang and  P.  Sun: Mater. Sci. Eng. A 764 (2019) 138244. doi:10.1016/j.msea.2019.138244
• 47)   L.  Wang,  H.  Ma,  Q.  Fan,  J.  Yao,  X.  Shen,  S.  Zhang,  Y.  Zhou,  Y.  Peng,  Y.  Gao and  D.  Wang: Mater. Charact. 176 (2021) 111112. doi:10.1016/j.matchar.2021.111112
• 48)   S.  Wang,  Z.  Hu,  W.  Zeng,  J.  Huang,  L.  Xiao and  H.  Zhou: Vacuum 191 (2021) 110404. doi:10.1016/j.vacuum.2021.110404
• 49)   Z.  Xu,  W.  Yang,  J.  Fan,  T.  Wu and  Z.  Gao: Materials 15 (2022) 3156. doi:10.3390/ma15093156
• 50)   K.  Takahashi and  E.  Sato: Tetsu-to-Hagané 98 (2012) 491–496. doi:10.2355/tetsutohagane.98.491
• 51)   M.  Abdel-Hady Gepreel and  M.  Niinomi: J. Mech. Behav. Biomed. Mater. 20 (2013) 407–415. doi:10.1016/j.jmbbm.2012.11.014
• 52)   A.  Yamamoto,  R.  Honma and  M.  Sumita: J. Biomed. Mater. Res. 39 (1998) 331–340. doi:10.1002/(SICI)1097-4636(199802)39:2%3C331::AID-JBM22%3E3.0.CO%3B2-E

Appendix

Fig. A1

XRD patterns of the Ti–xNb–yO alloys heat-treated under various conditions.