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Investigation of structure and mechanical properties: of Ti-7.2Al-2.9Mo--2.7W-3Nb-2.3Zr-0.4Si alloy.

Hardening of titanium alloys by solid solution alloying allows having additional possibilities due to influence on stability of high-temperature [beta]-phase and type and amount of products of its disintegration [1, 2].

Bi-phase ([alpha] + [beta])-alloys are characterized by good combination of strength and ductility and are the most numerous group of titanium alloys. However, it is difficult to achieve high strength of ([alpha] + [beta])-alloys at increased temperatures because of loss of strength stipulated by the phenomenon of structural super-ductility [3--6]. That's why working temperature of titanium ([alpha] + [beta])-alloys usually does not exceed 400-500 [degrees]C and problem of increasing their high-temperature strength remains actual. In this work influence of complex alloying on high-temperature strength and structural stability of bi-phase titanium ([alpha] + [beta])-alloy is considered.

Materials and methodology of investigations. Ingots of alloy for investigation were produced by double electron-beam cold hearth melting (EBCHM) and subjected to rolling in air according to standard technology for bi-phase alloys [7]. Content of alloying elements and impurities in the alloy was as follows, wt. %:
Base 7.2 2.9 2.7 3.0 2.3 0.4 0.18 0.026 0.0015

At the first stage rolling was started in b-area at temperature 1050 [degrees]C (temperature of polymorphous transformation is, approximately, 1020 [degrees]C). Five passes were made with general degree of deformation 55 %. At second stage produced billet of 32 [micro]m thickness was heated in ([alpha] + [beta])-area up to temperature 960 [degrees]C and subjected to rolling (11 passes) down to thickness 11 [micro]m. Temperature of the end of rolling was 850 [degrees]C, and general degree of deformation in ([alpha] + [beta])-area--more than 65 %.

Two modes of heat treatment were used for investigating structure of deformation after rolling: 850 [degrees]C, soaking for 2 h, cooling in air with subsequent ageing at 650 [degrees]C for 2 h; 950 [degrees]C, soaking for 2 h, cooling in air with subsequent ageing at 570 [degrees]C for 2 h. Hardness of the alloy in deformed state and after heat treatment was measured using hardness gage PMT-3 at the load 150 g, at which the imprint (approximately, 23 [micro]m diagonal) contained about 30 grains. Measurements were made using 10--12 imprints.

Temperature dependence of mechanical properties of the deformed alloy at static tension was defined within temperature range 20--750 [degrees]C. For this purpose flat specimens having 3.0x1.5 [micro]m section and length of working part 15 [micro]m were cut out from plates along the rolling. Tension was performed at rate of strain 1.2[10.sup.-3] [s.sup.1]. For estimating influence of atmosphere on the properties, specimens were tested in air and in vacuum.

For the purpose of determining temperature values, at which change of the type of elementary restructuring of atoms, which control plastic deformation, takes place, diagram in coordinates ln [[sigma].sub.02]-1/T was plotted, inclination of which was interpreted using expression [8]:

[[sigma].sub.02] = B exp ([DELTA]U/3kT), (1)

where dependence B upon temperature may be neglected [8, 9]; k is the Botlzmann's constant; T is the temperature, K; [DELTA]U is the change of internal energy connected with Gibbs activation energy; [DELTA]G is the characteristic of potential barrier, which is overcome when an elementary act of plastic deformation is performed according to the expression [DELTA]G = [DELTA]U--T[DELTA]S.


Investigation of the alloy structure was performed using transmission (TEM) and scanning (SEM) electron microscopes. SEM was also used for fractographic analysis of tested specimens.

Results of the experiment. Investigation of diffraction of electrons from the phases showed that the alloy mainly consisted of [alpha] and [beta]-phases (Figure 1). Identical in all investigated states of the alloy precipitations of particles of 0.3--0.4 [micro]m size were detected (Figures 1, 2). It was established in study of the electron diffraction that these were particles of silicide [Ti.sub.5][Si.sub.3], which, probably, formed in the process of the molten ingot cooling due to reduction of silicon solubility in titanium in [beta] [right arrow] [alpha] transformation.

Heat treatment according to the first mode practically did not change size of the grains (1--5 [micro]m) and quantitative ratio of -[alpha] and [beta]-phases. Just a certain rounding of [alpha]-grains was registered (Figure 3, a, b). Heating up to 950 [degrees]C according to the second version of heat treatment caused growth of grains of [alpha]-phase and increase of its amount, probably, because of recrystallization and redistribution of alloying elements (Figure 3, c). Due to higher temperature of heating (950 [degrees]C) activation of [beta] [right arrow] [alpha] disintegration was noted. Inside [alpha]-grains acicular precipitates of [alpha]-phase with cross size of needles about 0.04 [micro]m were more clearly seen (see Figure 2).

Microhardness of the alloy after rolling was (4.0 [+ or -] 0.3) GPa. Due to heat treatment according to the first mode its value increased up to (5.0 [+ or -] [+ or -] 0.28) GPa, but according to second one increased too but insignificantly--up to (4.4 [+ or -] 0.26) GPa. Evidently, after air hardening from 850 [degrees]C, heating up to 650 [degrees]C enabled disintegration of metastable b-phase, which caused additional hardening of the alloy. After heat treatment according to the second mode, loss of strength, which usually accompanies recrystallization (Figure 3, c), was compensated by disintegration of metastable [beta]-phase during repeated heating up to 570 [degrees]C and soaking at this temperature.


Temperature dependence of the alloy properties after rolling at static tension is presented in Figure 4. As test temperature increases from 20 to 650 [degrees]C, values of strength characteristics change insignificantly, and at further increase of temperature drastically reduce. Relative elongation S0 and reduction in area [[psi].sub.0] start to increase drastically at temperature 550 [degrees]C. Above 650 [degrees]C starts growth of uniform elongation [[delta].sub.u]. It proves stabilization of strain and difficulty in forming the neck (Figure 5). As far as curves of temperature dependence of yield strength and ultimate tensile strength practically coincide above 650 [degrees]C, strain hardening at these temperatures is absent (see Figure 4). Stability of plastic strain is ensured mainly due to appearance of rate dependence of flow stress [4, 5, 10]. In this case development of the neck is inhibited at the initial stage, because localization of strain in it is accompanied by increase of the strain rate and due to this increase of the flow stress is required. Appearance of rate dependence of flow stress is accompanied by beginning of its sharp temperature dependence.


In Figure 6 temperature dependence of yield strength is presented in semi-logarithmic coordinates. According to [8], sections of zigzag line determine temperature ranges of action of different strain mechanisms. In this very place near respective sections of zigzag line values of activation energies of thermally activated processes, obtained using equation (1), are presented.

In Figure 7 character of fracture of specimens at different temperatures is shown. In fracture of the specimens, which failed at 20 [degrees]C, mean value of facets equals 1.0--1.5 [micro]m and corresponds to size of the grains, taking into account their, approximately, 50 % deformation in neck of the specimen (Figures 1--3). As test temperature increases up to 550 [degrees]C, fracture gets pit character, whereby size of pits achieves 5 [micro]m. On their walls the relief was detected, which proves presence of internal structural elements, commeasurable with size of separate grains. As temperature increases up to 600 [degrees]C, size of pits continues to grow.

One usually connects presence of pits on fracture surface with failure over boundaries of originated in the process of deformation cells in particles of the second phase [9, 11]. In this case particles in the bottom of pits were not detected. Size of pits exceeds size of grains, and taking into account stability of the structure at heating up to 850 [degrees]C can not be connected with elements of structure of strain origin.


Evidently, as test temperature increases within 500--600 [degrees]C range, strength of boundaries approaches strength of the grain body. At 650 [degrees]C this process finishes by the fact that fracture over structural elements ceases in general and a specimen is stretched and acquires shape of a needle (Figure 5). Noted above stabilization of deformation hinders its localization not just at macrolevel (shape of a specimen), but also at the level of the alloy microstructure. Stretching into the needle proves absence of strain localization over section of a specimen in separate elements of the material structure with formation of an internal neck.

Experiments showed that atmosphere, in which tests are carried out (air or vacuum), practically does not affect strength characteristics of the alloy. Lower values have ductile characteristics in tests in air than in vacuum (see Figure 4). Reduction in area in test in air is connected with premature fracture initiated by defects on the surface (Figure 8). Usually pit frac ture starts in center of a specimen, where character of loading is more rigid.


In test in air defects in the form of cuts are formed on surface as a result of oxidation. They are concentrators of stresses and reason of origination of pit fracture from the surface.

Reduction in area in consequence of oxidation in test in air starts to be registered at the temperature 550 [degrees]C (Figure 4). Influence of premature fracture on level of relative elongation gets notable only at 600 [degrees]C. It is explained by low contribution of last stages of deformation in the neck into elongation at a lower temperature (Figure 5). So, atmosphere of test does not affect temperature dependence of uniform elongation, because this characteristic is determined at early stages of strain, when action of surface defects does not exert its influence yet.

Discussion of the results. Possibility of using equation of type (1) for analysis of activation parameters of strain mechanisms within different temperature ranges is considered in [8, 9, 12--14]. For separation of the Gibbs free activation energy [DELTA]G, which is characteristic of potential barrier for elementary deformation act and allows identifying these acts, it is necessary to introduce assumptions based on model approximations, upon which significantly depends the result to be obtained. That's why such data require for independent check, and only after this they may be used as substantiation of models used for performing calculations. And although it is impossible to determine specific strain mechanism on the basis ofjust tensile tests, calculation of internal energy of a crystal, which characterizes potential barrier and conditions of its overcoming and equals


allows separating temperature ranges, within which act strain mechanisms of similar nature. Indefiniteness, connected with unknown value of entropic term T([DELTA]S), is superimposed by significant difference of potential barriers for different groups of strain mechanisms.



Value of activation energy, produced using equation (1) [5], may be divided into groups (Table), depending upon strain mechanisms and temperature ranges of their action. At temperature T strain mechanism with potential barrier not exceeding 50kT may be activated due to heat fluctuations [8, 14], whereby the strain will be controlled by the mechanism, which has maximum activation energy among all possible mechanisms for this material [15].

Of interest is investigation of chemical and phase composition and the alloy structure influence on set Classification of strain mechanisms by activation energy levels of possible strain mechanisms and shift of the range of their action over temperature scale. Alloying, which shifts beginning of strain control by diffusion mechanisms into the area of higher temperatures, enables increase of high-temperature strength.

Carried out investigations showed that overcoming of spot obstacles by separate dislocations (activation energy is abut 0.18 eV) in the studied titanium alloy in deformed state (grain size is 1--5 [micro]m) controls strain up to relatively high temperatures (0.4[T.sub.melt]). This, evidently, is result of solid solution hardening of titanium by refractory elements.

From mechanisms, connected with diffusion, starts to act at comparatively low temperatures (0.4[T.sub.melt]) mechanism with activation energy 3.7 eV. So high value of activation energy, which is usually connected with grain-boundary glide, established rate of uniform elongation, drastic reduction of flow stress, and cessation of fracture over boundaries of grains are facts, which allow identifying mechanism of deformation that acts at the temperature from 650 [degrees]C and higher as inter-grain glide, on which structural super-ductility is based [2, 3].

In titanium ([alpha] + [beta])-alloys glide over grain boundaries at increased temperature is facilitated due to additional possibility of removing nonconformities, which occur over boundaries, due to [alpha] [right and left arrow] [beta] transformations. This, probably, also causes increased propensity of bi-phase titanium alloys to super-ductility. It is, evidently, possible to shift inter-grain glide into the areas of higher temperatures due to precipitation over [alpha]- and [beta]-phases of disperse, non-deformed particles, for example, silicides.

Prediction of prospects concerning use of silicide is based on possibility of influence by application of heat treatment on precipitation of secondary silicides allowing for high solubility of silicon in [beta]-phase and its drastic temperature dependence, and on precipitation of silicides during disintegration of [beta]-phase.


1. It is established that solid-solution hardening of titanium-base alloys is especially efficient due to possibility of [alpha] [right and left arrow] [beta] transformation control, which ensures production of disperse structure and precipitation of different metastable phases that may be effected by application of heat treatments.

2. It is shown that solution hardening for production of high-temperature titanium-base alloys is limited by phenomenon of glide over grain boundaries, which causes drastic loss of strength of the alloys at temperature above 600--650 [degrees]C. For increasing high-temperature strength of titanium alloys it is necessary to use disperse hardening, in particular, in the studied alloy it is advisable to increase content of silicon.

3. It is determined that alloying of titanium alloys for the purpose of high-temperature strength increase should simultaneously increase their heat resistance, because in the process of test in air as early as at 550 [degrees]C on surface of the metal appear defects connected with oxidation, which effect its mechanical properties.

[1.] Kolachev, B.A. (1976) Physical metals science of titanium. Moscow: Metallurgiya.

[2.] Kollings, E.V. (1988) Physical metals science of titanium alloys. Moscow: Metallurgiya.

[3.] Solonina, O.P., Glazunov, S.G. (1976) Heat-resistant titanium alloys. Moscow: Metallurgiya.

[4.] Grabsky, M.V. (1975) Structure superplasticity of metals. Moscow: Metallurgiya.

[5.] Gulyaev, A.P. (1982) Superplasticity of steel. Moscow: Metallurgiya.

[6.] Kajbyshev, O.A. (1984) Superplasticity of commercial alloys. Moscow: Metallurgiya.

[7.] Aleksandrov, V.K., Anoshkin, N.F., Bochvar, G.A. et al. (1979) Semi-finished products of titanium alloys. Moscow: Metallurgiya.

[8.] Borisenko, V.A. (1984) Hardness and strength of refractory materials at high temperatures. Kiev: Naukova Dumka.

[9.] Trefilov, V.I., Milman, Yu.V., Firstov, S.A. (1975) Physical principles of strength of refractory metals. Kiev: Naukova Dumka.

[10.] Shtejberg, A.M., Gajduchenya, V.F. (1996) Influence of strain rate cycling on elongation of steel and copper specimens. Metally, 3, 151--156.

[11.] Rybin, V.V., Vergazov, A.N., Solomko, Yu.M. (1978) Principles of intragrained fracture of metals with bcc-lattice. Fizika Metallov i Metallovedenie, 46(3), 582--596.

[12.] Gibbs, G.B. (1969) On the interpretation of experimental activation parameters for dislocation glide. Phil. Mag., 20(166), 867 -872.

[13.] Indenbom, V.L., Orlov, A.N. (1973) Introductory article. In: Thermally-activated processes in crystals. Moscow: Mir.

[14.] Dorn, D., Mout, D. (1966) Physical principles of creep. In: New materials and methods of material and alloy examination. Moscow: Metallurgiya.

[15.] Ivens, A., Roulings, R. (1973) Thermally-activated deformation of crystalline materials. In: Thermally-activated processes in crystals. Moscow: Mir.


(1) I.N. Frantsevich Institute of Materials Science Problems, NASU, Kiev, Ukraine (2) E.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine
Temperature range [DELTA][sigma] eV Characteristic of group

(0.1-0.3) [T.sub.melt] 0.04-0.30 Overcoming of spot
 obstacles by separate
(0.3-0.5) [T.sub.melt] 0.6-1.2 Break-off from Cottrell
 atmospheres from
 interstitial impurities
0.5 [T.sub.melt] 2.0-4.0 Mechanisms defined by
 diffusion of vacancies
 and substitutional
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Author:Firstov, S.A.; Zamkov, V.N.; Brodnikovsky, N.P.; Topolsky, V.F.; Kotko, A.V.
Publication:Advances in Electrometallurgy
Article Type:Report
Date:Apr 1, 2006
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