1. Introduction
The continuous striving of mankind to improve quality of life demands new materials for all fields of science and technology, especially so for medicine, for implants in particular. Previously used cobalt alloys and stainless steel, despite their good strength characteristics and corrosion resistance, contain elements toxic to the human body and have an elastic modulus many times higher than the elastic modulus of human bone tissue [1,2,3,4,5]. So, the non-magnetic alloy ASTM F75-2007 [6] used earlier as a material for bone implants has tensile strength σv up to 860 MPa, yield strength σ0.2 ~ 759 MPa, and Young’s modulus E ~ 220 GPa, and the stainless austenitic steel AISI 316L replacing it has σv ~ 485 MPa, σ0.2 ~ 170 MPa, and E ~ 210 GPa.
Mechanical Compatibility with Bone Tissue: One of the most important aspects in the development of bone implant materials is their mechanical compatibility with bone. For bone implants, both the stiffness (Young’s modulus) and the deformability of the material must be considered. Human bone has a certain modulus of elasticity that is much lower than that of traditional metallic materials such as titanium alloy. Young’s modulus E is a measure of a material’s stiffness that describes its ability to resist tension or compression. The implant material should have the E value as close as possible to its values for bone (up to 30 GPa) [7,8]. In this case, such a material will provide better mechanical biocompatibility between the implant and the bone in which it is fixed, thus avoiding excessive local overstretching of bones and joints and uneven load distribution. If the implant is too stiff compared to the surrounding bone, it can absorb most of the mechanical load, resulting in weakening and atrophy of the bone in contact with the implant. This phenomenon is called stress shielding. In such cases, the bone is not subjected to normal physiological stresses and its structure may weaken [9,10,11,12]. Resulting in the need for repeated surgical intervention with implant replacement [13].
In recent years, α + β and β titanium (BCC (body-centered cubic) lattice) alloys have become a key focus of research as a material for the next generation of joint endoprostheses due to their potentially promising set of properties. The most commonly used titanium alloys to date are α titanium (CP Grade 4) and α + β titanium alloy (Ti-6Al-4V, CP Grade 5), which have, respectively, σv ~ 550 MPa and ~ 925 MPa, σ0.2 ~ 483 MPa and ~ 900 MPa, and E ~ 102 and ~ 113 GPa. These are excessive values of elasticity in relation to bone tissue, which has E ~ 20–35 GPa [7,8]. In addition, alloying titanium with aluminum and vanadium creates long-term risks for the patient’s health, since the neurotoxic effect of excessive amounts of aluminum is known, contributing to the development of Parkinson’s and Alzheimer’s disease when aluminum accumulates in the brain [14,15]. Additionally, the effects of excess aluminum on male sexual function have been noted [16]. Systematic exposure to vanadium can lead to diarrhea, nausea, loss of appetite and weight loss, and impaired male fertility [17,18,19]. At the same time, modern studies of biocompatibility of niobium, zirconium, alloys with them, as well as their oxides, do not reveal any risks in long-term use as materials for implantation in the human body [20,21,22,23,24].
Superelastic materials have a lower modulus of elasticity and can distribute the load better, minimizing the risk of stress shielding effects. This allows for the bone to continue to function normally, helping to preserve and strengthen the bone around the implant. Superelasticity effect helps bone implant materials withstand cyclic loading (e.g., walking, running, lifting weights) without fatigue cracking. Superelastic materials can compensate for small changes in shape under load and return to their original state without permanent deformation, preventing the development of microcracks and increasing the life of the implant.
In this context, it is necessary to develop titanium alloys with high mechanical properties, low E value, and safe chemical elements. In particular, it is known that titanium alloys with metastable β-phase (β’, BCC (body-centered cubic) lattice), have the lowest E value, which can be obtained by selecting the alloy composition and heat treatment from the (α + β) phase, while avoiding the release of brittle and hard ω-phase. Ti β-type alloys were confirmed to have the lowest Young’s modulus. It should be noted that the majority of Ti β-type alloys with low Young’s modulus contain Nb as a β-stabilizing alloying element due to its moderate β-phase-stabilizing ability and biocompatibility [25,26,27,28,29,30,31,32].
In addition to the literature data, a computational method for estimating the initial phase composition by the molybdenum equivalent method [27] was used to evaluate the stability of the β-phase and optimize the alloy composition, as well as the experience of a previous study [33,34,35].
One of the most promising systems of such alloys is the Ti-Zr-Nb system. In recent years, there has been a growing interest in Ti-Zr-Nb and Ti-Nb-Zr alloys for medical applications. These materials have attracted attention due to their favorable combination of properties that make them promising for use as implantable devices such as endoprostheses. The inclusion of zirconium (Zr) and niobium (Nb) in the titanium alloy provides both the necessary durability of implants and maintains their low density and excellent corrosion resistance, low modulus of elasticity, and, in some combinations of elements, superelasticity, which is critical to ensure long-term functionality in the body environment [36,37,38,39,40].
The purpose of this study is to develop and comprehensively evaluate a titanium alloy containing zirconium and niobium for use as a material for bone implants. In the course of this work, the task was set to comprehensively analyze the influence of composition and thermomechanical processing on the structural and phase characteristics, as well as on the mechanical, cyclic, and superelastic properties of the alloy. Particular attention was paid to the study of the Young’s modulus as a key parameter that determines the mechanical compatibility of the material with human bone, which is important for improving durability and functional efficiency in implants. In this work, the physical and mechanical properties of Ti-(36-40)Zr-9Nb alloys (at.%) were investigated.
2. Materials and Research Methods
2.1. Selection of the Object of Research
Three alloys were investigated (at.%): Ti-36Zr-9Nb, Ti-38Zr-9Nb, and Ti-40Zr-9Nb. Titanium produced by company VSMPO-AVISMA Corporation (Russia) was used, Zirconium and Niobium produced by company Chepetskiy Mechanical Plant (Russia) were used. The Mo equivalent is an empirical parameter representing the contribution of the alloying elements to the stability of the β-phase relative to Mo [32].
To perform the calculation, it is necessary to convert atomic percentages into mass percentages:
Let us take the example of the Ti-36Zr-9Nb alloy (Table 1):
Taking into account the mass percentages and literature data on the contribution of the alloying elements, we calculate
[Mo] eq = (x(i) × C(Mo))/(C(i)), where C(i) and C(Mo) are the second critical concentrations of the alloying element and molybdenum, respectively [32].
When estimating the molybdenum equivalent of a complex alloy, the effect of different β-stabilizers is considered additive and the influence of α-stabilizers and neutral hardeners is neglected. Thus,
[Mo] eq%Nb/3,6 + %Zr/14.
The effect of Nb on β-phase stability is 1/3.6, i.e., niobium is 3.6 times weaker in stabilizing the beta-phase and zirconium is 14 times weaker than molybdenum.
Calculation of the molybdenum equivalent for 1-Ti-9Nb-36Zr, 2-Ti-9Nb-38Zr, and 3-Ti-9Nb-40Zr:
[Mo] eq = %Nb/3.6 + %Zr/14 = 12.38/3.6 + 48.63/14 = 6.91
[Mo] eq = b/3.6 + %Zr/14 = 12.22/3.6 + 50.68/14 = 7.01
[Mo] eq = %Nb/3.6 + %Zr/14 = 12.02/3.6 + 52.87/14 = 7.11
It can be concluded that all the alloys studied have an α+β phase based on the data obtained and the information in Table 2 in the annealed state; therefore, by means of thermomechanical treatment, namely, plastic deformation and quenching, alloys with unstable β-phase, low Young’s modulus, and superelasticity effect can be obtained from them.
2.2. Production and Processing of Research Samples
The alloy samples were melted in a copper water-cooled crystallizer in an argon arc furnace with a non-expendable tungsten electrode L200DI (Leybold-Heraeus, Hanau, Germany) at an argon pressure of 40 × 103 Pa. The following were used as charge materials: titanium (Ti-1 grade), zirconium (Zr-1 grade), and niobium (Hb-1 grade). The materials were placed in a copper water-cooled tray of a vacuum electric arc furnace as follows: titanium was placed at the bottom of the crystallizer wells, then zirconium, and niobium on top. The order of arrangement of initial metals corresponded to the melting temperature of each material, and melting was carried out by arc directed from top to bottom. Additionally, zirconium was placed in a separate tray to act as a getter. The furnace was vacuumized to a residual pressure of 1.33 Pa, filled with inert argon gas of high purity. At first, the zirconium getter is melted after induction into the inert gas chamber to remove possible oxygen impurities in the gas. Next, melting of the starting materials is performed until the formation of single ingots. Remelting was carried out 7 times with ingots turning over for better mixing of initial components.
The obtained ingots have a dendritic structure. Homogenizing annealing in a vacuum resistance furnace ESKVE-1.7.2.5/21 ShM13 (NITTIN, Moscow, Russia) at 1000 °C for 4 h was used to homogenize the structure and remove stresses. To protect against oxidation, annealing was performed in a vacuum environment at 27 × 10−4 Pa.
Primary deformation of ~18 mm thick castings was carried out by warm rolling at preheating to a temperature of 600 °C on a double-roller mill, DUO-300 (AO Istok ML, Fryazino, Moscow Oblast, Russia), with partial absolute compressions per pass: 1.5 mm down to billet thickness of 4.0 mm (heating every 2 passes, total 10 passes), then 1.0 mm down to billet thickness of 2.0 mm (heating every pass, total 2 passes), and a further 0.5 mm down to final billet thickness—1 mm (heating before the first pass at thickness of 2 mm, total 2 passes). The billets were heated before deformation in a muffle furnace KYLS 20.18.40/10 (Hans Beimler, Leipzig, Germany) to a temperature of 600 °C for 20 min before the first rolling and for 5 min during intermediate annealing. The plates were cold-rolled from a thickness of 1.5 mm. The effect of heat treatment on the structure and mechanical properties of Ti-(36-40)Zr-9Nb alloys was investigated on plate samples with a final thickness of 1 mm. Hardening of the samples was carried out after heating for 5 min at 600 °C in room temperature water. During rolling and quenching, a thin oxide layer formed on the surface of the samples. This layer was removed by mechanical grinding.
Preparation of samples for metallographic study was carried out by sequential grinding on a Buehler Phoenix 4000 machine (Lake Bluff, IL, USA). The microstructure was revealed and analyzed by etching in a weak solution of hydrofluoric and nitric acids with distilled water in the ratio of 10% HF: 10% HNO3: 80% H2O for 20–30 s. After, the microstructure was rinsed with distilled water, and residual water was removed by air flow. Light microscopy was carried out on Altami MET 5C (Altami, Saint Petersburg, Russia) microscope using a high-resolution video camera built into the device and Altami Studio 3.5 software.
2.3. Methods of Research
The direct measurement method was used to estimate the grain size: This method uses scale bars on the image and image processing software. A virtual ruler on the image is used to measure the average diameter of the grains. Considering that the shape of grains is different, several straight line measurements (width and length) were taken for elliptical or polyhedral grains and the average value is calculated.
X-ray diffractograms were obtained on an ARL X`TRA instrument (Thermo Fisher Scientific (Ecublens) sarl, Ecublens, Switzerland), in CuKα-radiation in parallel beam geometry. The instrument was calibrated to the NIST SRM-1976a [41] standard pattern, and the error in the position of the reflexes did not exceed 0.01° 2q. The crystal lattice parameter was refined by extrapolation to q = 90° by the Nelson–Riley method using Origin-2017 software. The quantitative content of crystalline phases was estimated by the method of corundum numbers. Before study, the surface of the samples was cleaned with emery paper (P4000) with subsequent washing in ethyl alcohol and distilled water.
To study microhardness, the Vickers method was used at a load of 200 g and a dwell time of 15 s on microhardness tester 401/402-MVD, manufactured by Instron Wolpert Wilson Instruments (Norwood, MA, USA). The microhardness of plates after rolling and quenching at 600 °C was investigated.
Static mechanical tests were performed on a universal testing machine INSTRON 3382 (Instron, Norwood, MA, USA) with a tensile speed of 1 mm/min using flat specimens of 28 × 6 × 1 mm and a working part size of 15 × 3 × 1 mm. The number of specimens for each alloy type and treatment type was 9 pieces. Processing of test results in determining the characteristics of mechanical properties was carried out in accordance with GOST 1497-84 using INSTRON Bluehill 2.0 software. The following characteristics were determined: conditional yield strength σ0.2, tensile strength σv, and relative elongation δ. Determination of E values was carried out in static tensile testing of the specimens using an extensometer.
Cyclic tests to determine the hysteresis behavior of the material were performed on an INSTRON 3382 universal testing machine. For each specimen, 10 cycles were performed; in the first part of the cycle, the specimen was stretched at a rate of 1 mm/min until the working area was elongated by 8%, and in the second part of the cycle, the specimen was compressed at a rate of 1 mm/min to the original size. The best hysteresis behavior can be used to evaluate the superplasticity effect of the alloys studied. Strain measurement was performed with an additional probe mounted directly on the specimen with two clamps along the edge of the specimen work area, so that the tensile value of the specimen work area was more accurate and did not include the strain of the gripping parts of the specimen, which occurs when using tensile values from the INSTRON 3382 frame-mounted probe. The specimen dimensions were 56 × 10 × 1 mm and the workspace size was 32 × 6 × 1 mm.
The fractography study was carried out on a TM4000 (Hitachi, Tokyo, Japan) scanning electron microscope. For study in raster mode, the samples were glued on a copper substrate using conductive carbon glue, and the accelerating voltage was 15 kV.
3. Results
3.1. Light Microscopy of Microstructures of Ti-(36-40)Zr-9Nb Alloys
Images of the microstructure after melting, rolling, and rolling and quenching are shown in Figure 1, Figure 2, and Figure 3, respectively.
After melting, the microstructure of ingots has a pronounced dendritic structure. This structure is typical of titanium alloys containing refractory elements that have not undergone additional heat treatment and is due to earlier crystallization of niobium and zirconium in relation to titanium at a rather slow cooling rate, which cannot be increased in the furnace used.
The microstructure of the plates in the longitudinal and cross-section of the plates is homogeneous after both rolling and quenching and consists of β grains of polyhedral shape. After rolling, the grain size in diameter is in the range of 30 and 100 µm, and the uniformity of the grains can be explained by the preheating of the billets and contact with the massive cold rolls of the rolling mill. The grain size after quenching is in the range of 50 and 170 µm. Grain growth is associated with the onset of recrystallization.
3.2. X-ray Phase Analysis (XRF) of Ti-(36-40)Zr-9Nb Alloys
The results of X-ray phase analysis of the composition of plates after rolling and after rolling and quenching are presented in Table 3 and Figure 4. The alloys after rolling consist only of metastable β-Ti. This is due to the influence of β-stabilizer Nb and neutral element Zr [42]. On its own, zirconium is a neutral hardener, but in the presence of other β-stabilizers, it also weakly stabilizes the β-phase. The phase composition of ingots does not change after quenching. This is because the β-phase formed during heating before rolling is fixed. After rolling and quenching, the alloys also consist of only β-phase. Quenching stabilizes the metastable high-temperature phase, promoting a uniform composition.
3.3. Microhardness of Ti-(36-40)Zr-9Nb Alloys
The results of plate microhardness analysis after rolling and quenching are presented in Table 4.
Microhardness of plates after quenching insignificantly decreases for all studied compositions. Quenching does not change the phase composition and very insignificant decrease in hardness is associated only with some increase in grain size due to partial recrystallization. In the plate Ti-38Zr-9Nb microhardness after quenching practically did not change, due to the small heating time for the beginning of recrystallization.
3.4. Mechanical Properties
The results of mechanical tests of Ti-(36-40)Zr-9Nb alloy are presented in Table 5. Typical tensile diagrams are shown in Figure 5, Figure 6 and Figure 7.
Based on the diagrams and the values obtained, it is noted that after rolling hardening led to an increase in the ductility of all studied alloys Ti-(36-40)Zr-9Nb with an approximate preservation of the strength level. The alloy Ti-38Zr-9Nb showed the highest strength.
The linearly increasing section on each diagram demonstrates elastic deformation proceeding according to Hooke’s law. After it, a kinked area appears in the hardened samples, indicating the yield strength associated with reversible phase change: β’<-> α’’. The β`-phase has a crystal lattice like the β-phase, and the presence of a kink in the mechanical studies’ images can distinguish the state of the material in which the metastable β-phase is present. The α’’-phase is an intermediate phase of the β → α’-transformation and is the result of an incomplete shift and requires much smaller atomic displacements than the formation of the α- and α’-phase [43]. The phase transition is realized by a shear mechanism. The cause of the phase transition is loading and the plastic deformation that takes place. The occurring phase change means the manifestation of hysteresis behavior in the compositions. Following this site, a region is observed in which the load is practically unchanged and the specimen undergoes tensile strain.
3.5. Cyclical Loading
Figure 8, Figure 9 and Figure 10 show the tensile diagrams for cyclic loading up to 8% strain for 10 cycles.
Three compositions were studied at cyclic strain up to 8%, and the lowest residual strain was shown by alloy 38%; therefore, additional studies were made on this alloy at cyclic strains of 2, 3, and 6%. It was of interesting to observe at what maximum strain the residual strain would be as close to 0 as possible. For Ti-38Zr-9Nb, the samples were cyclically stretched to 2%, 3%, 6%, and 8% (Figure 11).
Ti-36Zr-9Nb exhibited a residual strain equal to 1.5%.
Ti-38Zr-9Nb returns with the lowest residual strain when loaded to 2% and 3%. At 6%, the residual strain is 0.4%, and at a loading of 8% = 0.7%.
Ti-40Zr-9Nb exhibits a residual strain equal to 2.4%, which is the worst result.
The linear plot in each diagram demonstrates the transition of metastable β-phase to α’’. In this regard, it can be concluded that the best hysteresis behavior, which characterizes the manifestation of the superelasticity effect, shows the Ti-38Zr-9Nb alloy.
3.6. Young’s Modulus
The results of the study of the elastic modulus of plates after rolling and quenching are presented in Table 6.
In general, the E values of all alloys were approximately the same and are already quite close to the E value for human bone tissue. According to the test results, the alloy Ti-40Zr-9Nb, compared to the others, had slightly lower E values, which is due to the largest amount of Zr, which weakens the bond between atoms in the presence of β-stabilizer.
3.7. Fractography
Figure 12 shows a typical view of the fracture surface under static loading for samples of Ti-40Zr-9Nb alloy after rolling and quenching: 1—crack nucleation surface; 2—ductile fracture surface; 3—brittle fracture surface before fracture. Fracture direction from surface 1 to the center of the specimen.
In the fracture zones, there are both brittle fracture areas in the form of “jet patterns” and ductile areas in the form of “pits”. When the specimens are hardened, the brittle fracture areas decrease; more ductile fracture areas appear; the crystallographic character of the fracture takes the form of shearing or bilateral shearing to the center of the specimen; the appearance of fractures becomes more fibrous or ductile, which corresponds to the results of mechanical tests; and the formation of a neck is observed. In all specimens, the crack leading to failure occurred on the surface and propagated through the center to the opposite outside of the specimens, or from two opposite surfaces to the center.
4. Conclusions
It was found that alloys (at. %) Ti-36Zr-9Nb, Ti-38Zr-9Nb, and Ti-40Zr-9Nb both after rolling and after rolling and quenching have a homogeneous microstructure consisting of polyhedral grains of β-Ti stabilized by Nb and Zr. The grain size after rolling is ~100 μm and after quenching is ~170 μm.
Quenching is rational for the studied alloys, as it has practically no effect on Young’s modulus E, microhardness, or strength characteristics, but increases plasticity and the manifestation of the superelasticity effect. The achieved values of E ~ 50–60 GPa are quite close to the E values in human bone tissues.
Alloy Ti-38Zr-9Nb has the most balanced complex of mechanical characteristics in comparison with the other two alloys: at approximately the same values of E and microhardness, it has greater strength and plasticity characteristics, and it has a more pronounced effect of superelasticity.
In connection with the rather high mechanical properties (ultimate strength σv ~ 630–660 MPa, plasticity δ ~ 14–23%) of the studied alloys and more optimal modulus of elasticity in relation to pure titanium and titanium Grade 5, it seems appropriate to further study the material for biological compatibility and corrosion resistance.
Conceptualization, S.V.K. and L.A.S.; Data curation, M.A.K., E.O.N. and L.A.S.; Formal analysis, K.V.S. and S.V.K.; Funding acquisition, M.A.S. (Mikhail A. Sevostyanov) and A.G.K.; Investigation, K.V.S., S.V.K., M.A.K., Y.G., M.A.S. (Maria A. Sudarchikova), T.M.S. and S.A.M.; Methodology, S.V.K., A.D.G., M.A.S. (Maria A. Sudarchikova), T.M.S. and S.A.M.; Project administration, M.A.S. (Mikhail A. Sevostyanov) and A.G.K.; Resources, E.O.N. and A.G.K.; Supervision, A.D.G. and A.G.K.; Validation, A.D.G. and Y.A.M.; Visualization, Y.G. and Y.A.M.; Writing—original draft, K.V.S. and Y.G.; Writing—review and editing, M.A.K. and M.A.S. (Maria A. Sudarchikova) All authors have read and agreed to the published version of the manuscript.
Not applicable.
The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author(s).
The authors declare no conflicts of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript; or in the decision to publish the results. All the authors have reviewed the manuscript and agree on submission to your journal.
Footnotes
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Enlarge this image.Figure 1. Microstructure of the ingots after melting: (a)—Ti-36Zr-9Nb; (b)—Ti-38Zr-9Nb; (c)—Ti-40Zr-9Nb.
Enlarge this image.Figure 2. Microstructure of the plates after rolling: (a)—Ti-36Zr-9Nb; (b)—Ti-38Zr-9Nb; (c)—Ti-40Zr-9Nb.
Enlarge this image.Figure 3. Microstructure of the plates after rolling and quenching: (a)—Ti-36Zr-9Nb; (b)—Ti-38Zr-9Nb; (c)—Ti-40Zr-9Nb.
Enlarge this image.Figure 4. X-ray diffraction patterns of alloys: (a)—Ti-36Zr-9Nb; (b)—Ti-38Zr-9Nb; (c)—Ti-40Zr-9Nb.
Enlarge this image.Figure 4. X-ray diffraction patterns of alloys: (a)—Ti-36Zr-9Nb; (b)—Ti-38Zr-9Nb; (c)—Ti-40Zr-9Nb.
Enlarge this image.Figure 12. Fractography of Ti-40Zr-9Nb alloy after rolling and quenching: (a) general view of static fracture; (b) view of fracture surface.
Conversion of atomic percentage to mass percentage.
| Chemical Element | Atomic Mass (a.e.m.) | Atomic Percentage (%) | Atomic Mass * Atomic Percent = | Mass Percentage (%) |
|---|---|---|---|---|
| Ti | 47.87 | 55 | 2632.85 | 38.99 |
| Nb | 92.91 | 9 | 836.16 | 12.38 |
| Zr | 91.22 | 36 | 3283.92 | 48.63 |
Value of molybdenum equivalent as a function of alloy phase composition (Adapted from Ref. [
| Alloy Grade | [Mo] Equivalent |
|---|---|
| (α + β)-alloys | 3.3–10 |
| Pseudo β-alloys | 11.0–16.3 |
| β-alloys | 16.9 |
Phase composition of Ti-(36-40)Zr-9Nb alloy plates before and after quenching.
| Composition | Condition | Phase Structure | Type of Crystal Lattice | Contents, Volume. % | Crystal Lattice Parameters |
|---|---|---|---|---|---|
| Ti-36Zr-9Nb | After the rolling process | β-Ti | type A2, cI2 | 100 | a = 3.40141 ± 0.00006 Å |
| After the rolling process and quenching | β-Ti | type A2, cI2 | 100 | a = 3.40210 ± 0.00002 Å | |
| Ti-38Zr-9Nb | After the rolling process | β-Ti | type A2, cI2 | 100 | a = 3.40842 ± 0.00005 Å |
| After the rolling process and quenching | β-Ti | type A2, cI2 | 100 | a = 3.40913 ± 0.00004 Å | |
| Ti-40Zr-9Nb | After the rolling process | β-Ti | type A2, cI2 | 100 | a = 3.42241 Å |
| After the rolling process and quenching | β-Ti | type A2, cI2 | 100 | a = 3.41362 ± 0.00004 Å |
Microhardness of Ti-(36-40)Zr-9Nb plates.
| Composition | Condition | Microhardness, HV |
|---|---|---|
| Ti-36Zr-9Nb | After the rolling process | 223 ± 2 |
| After the rolling process and quenching | 219 ± 1 | |
| Ti-38Zr-9Nb | After the rolling process | 216 ± 2 |
| After the rolling process and quenching | 216 ± 2 | |
| Ti-40Zr-9Nb | After the rolling process | 219 ± 2 |
| After the rolling process and quenching | 213 ± 2 |
Mechanical properties of Ti-36Zr-9Nb plates.
| State | δ, % | σ0.2, MPa | σB, MPa |
|---|---|---|---|
| Ti-36Zr-9Nb | |||
| After the rolling process | 11.8 ± 1.9 | 484 ± 69 | 637 ± 21 |
| After the rolling process and quenching | 19.3 ± 2.0 | 289 ± 20 | 628 ± 20 |
| Ti-38Zr-9Nb | |||
| After the rolling process | 14.4 ± 2.1 | 494 ± 35 | 656 ± 6 |
| After the rolling process and quenching | 23.0 ± 1.8 | 324 ± 17 | 613 ± 15 |
| Ti-40Zr-9Nb | |||
| After the rolling process | 8.9 ± 1.0 | 532 ± 23 | 604 ± 15 |
| After the rolling process and quenching | 23.4 ± 0.5 | 364 ± 25 | 594 ± 29 |
Young’s modulus values of plates.
| Composition | E, GPa |
|---|---|
| Ti-36Zr-9Nb | 55 ± 2 |
| Ti-38Zr-9Nb | 59 ± 3 |
| Ti-40Zr-9Nb | 51 ± 2 |
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Details
; Kaplan, Mikhail A 1
; Gorbenko, Artem D 1
; Guo, Yucheng 2 ; Nasakina, Elena O 1
; Sudarchikova, Maria A 1 ; Sevostyanova, Tatiana M 3 ; Morozova, Yaroslava A 1 ; Shatova, Lyudmila A 4 ; Mikhlik, Sofia A 1 ; Sevostyanov, Mikhail A 1 ; Kolmakov, Alexey G 1 1 A.A. Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences (IMET RAS), 119334 Moscow, Russia;
2 Department of Materials Science, Bauman Moscow State Technical University (BMSTU), 105005 Moscow, Russia;
3 National Medical and Surgical Center Named after N.I. Pirogov of the Ministry of Health of the Russian Federation, 117513 Moscow, Russia;
4 Faculty of Radio Engineering and Electronics, Voronezh State Technical University, 394026 Voronezh, Russia;