- Email: [email protected]
Effect of hydrogen on internal friction and elastic modulus in titanium alloys
Available online at www.sciencedirect.com
Acta Materialia 57 (2009) 715–721 www.elsevier.com/locate/actamat
Effect of hydrogen on internal friction and elastic modulus in titanium alloys N.L. Arabajian a, V.I. Serdobintsev a, V.M. Tavkhelidze a,*, T.A. Peradze b, Yu.I. Stamateli b, K.M. Gorgadze b a
E. Andronikashvili Institute of Physics, 6 Tamarashvili St., 0177 Tbilisi, Georgia b Georgian Technical University, 77 Kostava St., 0175 Tbilisi, Georgia
Received 20 July 2008; received in revised form 4 October 2008; accepted 9 October 2008 Available online 10 November 2008
Abstract The effect of hydrogen on the variation with temperature of internal friction (QI) and elastic modulus (E) of a number of Ti-based alloys has been studied in the Hz and kHz frequency ranges. A relaxation peak of internal friction with a high degree of relaxation (QImax 101) and with a DE effect is observed in all hydrogen-doped samples at T 600 K at 1 kHz, and at T 500 K at 1 Hz. Such a peak is not present in samples without hydrogen. The activation energy W and the frequency factor v0 of the observed relaxation are determined to be W 1.55 eV, v0 1017 s1. It is shown that the observed effects are connected with the mechanism of grain boundary relaxation, as the introduction of hydrogen into titanium alloys leads to the formation of fine-grained structures. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Titanium alloys; Internal friction; Hydrogen; Grain boundaries relaxation
1. Introduction Titanium alloys with b-isomorphic elements (Nb, V, Ta, Mo, Zr) are structural materials that exhibit a shape memory effect (SME) and are used in different spheres of engineering. In contrast to the well-known and widely studied intermetallic alloys Ti–Ni, the materials investigated in this work are substitutional solid solutions in which the doping elements are disordered. In these alloys, the SME is based on thermoelastic phase transformation of body-centered cubic b-austenite into orthorhombic a-martensite [1]. At present, the following effects, which are connected with the introduction of hydrogen into titanium alloys, are known: a) decrease in the temperature of the martensite transition [2];
b) changes in the energy of stacking faults and in the energy of interphase and intergrain boundaries [3]; c) hydrogen-phase cold hardening [4]. Measurements of internal friction and elasticity modulus have shown that hydrogen exerts a significant influence on the elastic characteristics and damping ability of titanium alloys [2,5–7]. Although these effects connected with hydrogen are significant, their structural mechanisms are not yet quite clear. Thus, further investigation of the properties of hydrogenous Ti alloys is of a special interest. The present work deals with the investigation of the effect of hydrogen on the characteristics of the martensite transformation, elastic properties and damping ability of Ti alloys with Nb, V, Ta, Mo, Zr, over a wide range of temperatures and hydrogen concentrations. 2. Materials and methods
*
Corresponding author. E-mail addresses: [email protected], [email protected] (V.M. Tavkhe lidze).
Measurements were carried out on the alloys: Ti– 50.1Ta, Ti–59.8Ta, Ti–33Nb–7Zr, Ti–5.1Ta–4.9Mo–4.9 V,
1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.10.013
716
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
Ti–5.1Ta–4.9Mo–4.9 V–2.0Zr (all wt.%), obtained by multiple electron-beam remelting with subsequent hot rolling. The measurements of internal friction (QI) and elastic modulus (E) were carried out by the method of excitation of bending vibrations of the sample in the kHz frequency interval and torsional vibrations on a Coulomb-type pendulum in the frequency range 1–4 Hz. The internal friction was determined by the damping decrement of natural vibrations of the sample by means of an electronic discriminator registering the number of vibrations at half the vibration amplitude, and the elastic modulus was determined by the frequency f of resonance vibrations (according to the relation E / f 2 . The samples were thin plates 0.5 1 20 mm3. The initial phase state of the alloys was obtained by hardening in water after isothermal exposure in the temperature range (T > 1100 K) at which high-temperature b-austenite exists. Depending on the type of doping elements and their concentrations, the initial state of the alloy corresponded to (b00 +a00 )- or a00 -phase [8]. The characteristic temperatures of direct (Ms, start; Mf, finish) and inverse (As, start; Af, finish) martensite transformations were determined by differential thermal analysis (DTA) and from the curves of deformation – temperature dependence for torsional deformation under the action of permanent loading in the process of cooling the samples from a certain temperature above the martensite transformation and a subsequent heating. This method is applicable for determining the temperatures of the start and the finish of the martensite transformation owing to the SME. Introduction of hydrogen into samples was carried out in a special high-vacuum chamber at a temperature of 800 °C. After the chamber was filled with gaseous hydrogen, the change in hydrogen pressure was traced. Further, when the process of saturation of the sample with hydrogen was ended, the furnace in which the chamber with the sample was placed was rapidly removed, and the sample was cooled to room temperature in 10 s. The hydrogen concentration in the sample was determined by the pressure drop in the chamber during the hydrogenation process. 3. Experimental results 3.1. Ti–59.8 wt.% Ta alloy Fig. 1 shows the temperature spectra of internal friction and elastic modulus (in the units of sample frequency squared) of Ti–59.8 wt.% Ta alloy measured on a torsion pendulum with heating at the rate of 8 K min1. The form of the QI and f 2 curves is typical for phase transitions of the martensite type. As known [9], the internal friction at a martensite transformation is described by the relation:
Fig. 1. Temperature dependencies of (a) f2 and (b) QI of Ti–59.8 wt.% Ta alloy measured on a torsion pendulum during heating.
QI ðT Þ ¼
1 nE dT du ; 2 ra x dt dT
ð1Þ
where n is the shear deformation of the lattice at its structural rearrangement, E is the elasticity modulus, ra is the alternating voltage amplitude, T is the temperature, t is the time, x is the vibration frequency and u is the volume fraction of martensite phase. The measurements in the kHz frequency range give similar temperature dependencies for QI and f 2 at the excitation of the flexural mode. Fig. 2 (curve 1) shows the dependence QI (T) of the same sample measured at a frequency of 1 kHz. The sample heating rate was again 8 K min1. Comparison of Figs. 1 and 2 shows that the amplitude of the QI peak at the frequency of 4 Hz is about an order of magnitude higher than at the frequency of 1 kHz, though it does not fit Eq. (1). To check the dependence of peak height QI on the heating rate dT/dt, the experiment was carried out with a sharp change of T (curve 2 in Fig. 2). At TI = 480 K, the sample heater was switched off near the top of the peak, causing a sharp decrease in the heating rate, which drops to zero. As Fig. 2 shows, after stopping the heating, the increase in the internal friction sharply slows down and then decreases almost to the background value (the time of decrease is 15 min). The repeated heating at T2 = 500 K shows that QI begins to increase rapidly and practically joins curve 1 in Fig. 2. This observation proves the fact that the peak registered at the temperature of 500 K is indeed connected with the mar-
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
717
Fig. 2. Temperature dependencies of QI of Ti–59.8 wt.% Ta alloy measured in a flexural mode: curve 1, heating rate 8 K min1; curve 2, measurements in which heating was stopped at T1 = 480 K and resumed at T2 = 500 K.
tensite–austenite phase transition. The measurements by DTA give the following temperatures for the start and finish of the transition: As = 450 K, Af = 520 K. 3.2. Ti–59.8 wt.% Ta alloy doped with hydrogen After investigation of the initial samples of Ti–59.8 wt.% Ta alloy, hydrogenation was carried out according to the method described in Section 2. Three samples were prepared with hydrogen concentrations 0.07, 0.10 and 0.17 wt.% H. Fig. 3 shows the temperature dependencies of QI and 2 f of the Ti–59.8 wt.% Ta + 0.07 wt.% H sample measured during the process of heating at the excitation of the flexural mode. (It should be noted that the specific character of the measurement of QI using an electron discriminator does not unfortunately allow measurements of the values of QI 0.1 with an error <25%.) The shape of the curves points to the existence of a relaxation process with an unusually high degree of relaxation. The absence of the peak connected with martensite–austenite phase transition is also noteworthy, although the DTA measurements and the estimation of transition temperature by the method of sample bending show that the phase transition takes place in the temperature range 395–495 K. To estimate the activation energy of the relaxation process, the measurements of temperature dependence of QI and f 2 of the same sample were carried out on the torsion pendulum (Fig. 4). As seen from this figure, the temperature of the relaxation maximum almost coincides with the temperature of phase transition, and therefore it was checked whether the observed peak is caused only by relaxation or by the superposition of two processes: relaxation and phase transition. For this purpose, at the peak maximum, the temperature was stabilized and kept for 1 h; however, no changes in the level of vibration energy absorption were observed during this period.
Fig. 3. Temperature dependencies of (a) f2 and (b) QI of Ti–59.8 wt.% Ta + 0.07 wt.% H sample measured in a flexural mode in process of heating (As = 395 K, Af = 495 K).
Fig. 4. Temperature dependence of QI of Ti–59.8 wt.% Ta + 0.07 wt.% H sample measured on a torsion pendulum (j heating, h cooling).
In addition, a comparison of the QI (T) curves recorded during heating and cooling of the sample shows the absence of hysteresis in phase transitions of the first kind, which also indicates the existence of only a relaxation process. The estimation of activation energy of the process according to the temperature shift of the relaxation maximum in accordance with the Arrhenius equation gives the following values of activation energy W and frequency factor v0: W = 1.56 ± 0.02 eV, v0 = 1017 ± 0.5 s1. The measurements of QI (T) of Ti–Ta samples with 0.1 and 0.17 wt.% H show that the temperature position of the
718
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
observed relaxation peak does not depend on hydrogen concentration, and the degree of relaxation changes insignificantly (QImax 0.1 for 0.1 and 0.17 wt.% H concentration). 3.3. Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V alloy Fig. 5 shows QI(T) and f 2(T) dependencies of Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V alloy. The trend of the curves shows that in the temperature range 290– 770 K, any anomalies that can be connected either with the phase transition or with the relaxation are not observed, although the measurements of heat capacity (DTA) and SME show the martensite–austenite phase transition in the range of 670–750 K. One can assume that the peak of QI corresponding to the martensite–austenite transition is masked by the rapidly increasing background of internal friction. 3.4. Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H alloy Fig. 6 shows QI(T) and f 2(T) dependencies of Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H alloy measured at 1 kHz frequency. As is seen from the figure, the introduction of hydrogen into the sample changes the essential form of temperature spectra of QI and f 2.
Fig. 6. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H sample measured in a flexural mode in process of heating (As = 330 K, Af = 470 K).
As in the case of Ti–Ta + H system, in the range of T 610 K there appears a relaxation peak with a very high degree of relaxation (QImax 0.1) with a corresponding anomaly of the elastic modulus. The measurements at a low frequency (Fig. 7) allow us to calculate the activation energy of the peak W and the frequency factor v0 by the temperature shift of QImax. The obtained data W = 1.53 ± 0.02 eV and v0 = 1016±0.5 are in good agreement with the corresponding data for Ti–Ta + H system, which means that the nature of the observed peaks in both systems can be assumed to be identical. Measurements of the start (As) and finish (Af) temperatures of the martensite transition give values of As = 330 K and Af = 470 K. Thus, the introduction of hydrogen into the sample decreased the temperature of the martensite–austenite transition by about 200 K. This influence of hydrogen on the temperature of the martensite transformation was found earlier [2] in a study of industrial Ti-based alloys, and is explained by a stabilizing action of hydrogen on the high-temperature b-phase of the alloy. As in the case of the Ti–Ta + H system, the presence of the phase transition is not observed on the curve of QI(T) dependence. 3.5. Ti alloys of different compositions doped with hydrogen Fig. 5. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V sample measured in torsional vibration mode during heating (As = 670 K, Af = 750 K).
Apart from the measurements described above in detail, the measurements of the following alloys were also carried out: Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V–2.0 wt.%
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
Fig. 7. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H sample measured in torsional vibration mode.
719
placed in a furnace heated to 800 °C to provide a rapid heating of the sample. The observation of the dynamics of pressure in the chamber has shown that within 10 min the sample is completely dehydrogenated. After reaching the ultimate vacuum (104 Pa), the furnace was quickly removed from the chamber, and the sample was cooled down to room temperature in 10 s. Fig. 8 shows the temperature spectra of QI (curve 1) and f 2 of the sample after degassing recorded at the flexural mode during the heating process. From curve 1 it is clear that the relaxation peak observed in the alloy with hydrogen is present again in the spectrum; however, its width is somewhat larger, and the temperature position of the maximum is shifted to lower temperatures. In addition, the behavior of the elastic modulus is not typical of a simple relaxation process, indicating, in our opinion, the reconstruction of internal structure of the sample. After heating up to 800 K during the measurement process and subsequent slow cooling, the measurements of QI (T) and f 2(T) were carried out again. The results are given in Fig. 8 (curve 2). The temperature position of the QI(T) peak, its amplitude and shape show a significant difference from the spectrum in curve 1. Checking the origin of the QI peak by stopping the heating near the temperature of the maximum has shown that the peak is caused by the martensite–austenite phase transition (Fig. 8, curve 3).
Zr + 0.3 wt.% H; Ti–33 wt.% Nb–7 wt.% Zr + 0.16 wt.% H; Ti–50.1 wt.% Ta + 0.21 wt.% H. The measurements have shown that after the introduction of hydrogen, in all these systems there appears a relaxation peak of internal friction of a high magnitude with the maximum at T 610 K at 1 kHz and at T 510 K at 1 Hz; hence it can be assumed that the mechanism of relaxation is the same for all systems studied here. 4. Discussion It is absolutely clear that the processes connected with the diffusion of interstitial hydrogen cannot be characterized by such high values of activation energy obtained in our experiments. For example, according to data [7] obtained for Ti– Ni–Cu alloy, the activation energy of relaxation peak connected with the motion of hydrogen is 0.6 eV. The same can be said about the relaxation processes of Snoek– Ko¨ster-type, which are connected with the ordering of hydrogen atoms considered as elastic dipoles in dislocation field, as well as about the processes of reorientation of substitutional atom–interstitial atom pairs [10]. To check the reversibility of the observed phenomena connected with the introduction of hydrogen into alloy, the studied Ti–59.8 wt.% Ta + 0.07 wt.% H sample was placed into a vacuum chamber for dehydrogenization. After reaching a vacuum of 104 Pa, the chamber was
Fig. 8. Temperature dependencies of (a) f2 and (b) QI of Ti–59.8 wt.% Ta sample after the dehydrogenation measured in a flexural mode. Curve 1, the first measurement during heating to T = 800 K; curve 2, the second measurement after heating and subsequent slow cooling; curve 3, measurement when the heating was stopped at T 500 K.
720
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
DTA and SME measurements confirm our statement. Just the same results are obtained by measuring the samples with 0.10 and 0.17 wt.% H after dehydrogenation. Hence, it can be assumed that the relaxation appearing after the introduction of hydrogen is connected with it indirectly; it is the result of the change of microstructure of the sample after hydrogenation. To explain such unusual behavior of the alloy, it is necessary to use the investigations carried out in Ref. [3] on Ti alloys with Al, which are similar in their properties to the alloys studied in the present work. In Ref. [3] it was observed that the introduction of hydrogen into samples by a method similar to ours leads to the appearance of a fine-grained structure. It is therefore logical to connect the effect discovered by us to grain boundary relaxation [10]. Although the existing mathematical models cannot describe exactly the grain boundary peaks of internal friction observed in many metals and alloys, the available experimental data [11–14] testify that in most cases the activation energy is close to 1.5 eV, and the degree of relaxation is of the same order (QImax 0.1) as in our experiments. Furthermore, according to the theory, the grain boundary peak is observed only when the grain size is less than the thickness of the sample (in our case 0.5 mm). It is also known that high-temperature annealing leads to an increase in the grain size (crystallites), and therefore the disappearance of the peak at T 600 K after heating up to T 800 K can be explained by the increase in the grain size up to >0.5 mm. The assumption of the formation of a fine-grained structure in the process of hydrogenation also explains the absence of the Q1(T) peak corresponding to martensite phase transition. Indeed, according to Eq. (1), the internal friction is proportional to the macroscopic shear strain n of the sample. On the other hand, as a result of hot rolling, the initial samples have a pronounced texture, owing to which the formed martensite plates (lamellas) are oriented predominantly in one definite direction. Therefore, the macroscopic shear strain makes a contribution to damping. As for the formation of a fine-grained structure as a result of hydrogenation, this causes isotropic arrangement of crystallites containing the martensite phase. Hence, the macroscopic shear strain does not arise as a result of martensite transition, and therefore the transition is not observed in Q1(T) spectra. After dehydrogenation and high-temperature annealing, the anisotropy in the orientation of martensite plates again becomes significant and leads to the appearance of the Q1 peak in the temperature range of martensite transformation. The dehydrogenation of Ti–5.1 wt.% Ta–4.9 wt.% Mo– 4.9 wt.% V + 0.43 wt.% H sample gives a result slightly different from that described above. Namely, after dehydrogenation at 800 °C and a rapid cooling down to room temperature, Q1(T) peak (Fig. 9) corresponding to grain-boundary relaxation decreases manifold and is almost indiscernible even at the first heating. This can be explained by a different kinetics of the increase of the grain size of Ti – 59.8 wt.% Ta and Ti–5.1 wt.% Ta–4.9 wt.%
Fig. 9. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V sample measured in flexural mode immediately after dehydrogenation.
Mo–4.9 wt.% V systems strongly differing from each other in the degree of doping by b-isomorphic elements. Thus, it can be stated that the hydrogen introduced into the studied alloys destroys the textured structure of martensite lamellae, leading to the appearance of a fine-grained structure with a high concentration of grain boundaries. Therefore, after introduction of hydrogen into the sample, the internal friction peak connected with the martensite transformation disappears, and a new relaxation peak appears due to the motion of grain boundaries. After the hydrogen is removed from the sample, the relaxation peak is still observed until the subsequent annealing restores the large-block structure of martensite lamellae. 5. Conclusions 1. It is shown that the introduction of hydrogen into the investigated Ti alloys causes the appearance of a Debye peak of internal friction with a high relaxation degree (Q1max 101), and the peak connected with the martensite phase transformation disappears. 2. The parameters of the discovered relaxation process, activation energy W and frequency factor v0, indicate that the relaxation cannot be connected with the process of hydrogen atoms jumping between interstitial positions in the sublattice. 3. It is shown that the discovered relaxation is indirectly connected with hydrogenation of samples and can be
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
explained by the processes of transformation of the crystal lattice from a large-block-textured state into an isotropic fine-grained state. 4. The analysis of the obtained data leads to the conclusion that the observed QI peak is caused by the mechanism of grain boundary relaxation, which is the result of viscous slipping of neighboring grains with respect to each other. 5. The impossibility of observing the martensite peak of internal friction in hydrogenated samples is also connected with the formation of a fine-grained structure and, as a consequence, with the absence of macroscopic shear strain in the process of martensite transformation. Acknowledgments The authors thank Dr. V. Melik-Shakhnazarov for his valuable comments and N. Goldbaum for her technical assistance in preparing the manuscript.
721
References [1] Zvicker U. Titan and titanlegierungen. Berlin, Heidelberg, New York: Springer-Verlag; 1974. [2] Ilyin A, Kollerov M, Golovin I. J Alloys Comp 1997;253:144. [3] Kolachev B, Nosov V. FMM 1984;57(2):286 [in Russian]. [4] Goltsov V. Mater Sci Eng 1981;49(2):109. [5] Mazzolai FM, Biskarini A, Campanella R, Collusi B, Mazzolai G, Rotini A, et al. Acta Mater 2003;51:573. [6] Biskarini A, Collusi B, Mazzolai G, Tuissi A, Mazzolai FM. J Alloys Comp 2003;355:52. [7] Fan G, Zhou Y, Otsuka K, Ren X, Nakamura K, Ohba T, et al. Acta Mater 2006;54:5221. [8] Lyasotskaya V, Lyasotskyi I, Ravdonikas I, Belov S. In: Proceedings of institutes of higher education, vol. 3, 1985. p. 82. [9] Malygin GA. Physics-Uspekhi 2001;44(2):173. [10] Novick AS, Berry BS. Anelastic relaxation in crystalline solids. New York: Academic Press; 1972. [11] Ke T. Phys Rev 1947;71:533. [12] Weinig S, Mashlin E. Trans Met Soc AIME 1957;209:32. [13] Cordea J, Spretnak J. Trans Met Soc AIME 1966;236:1685. [14] Williams T, Leak G. Acta Met 1967;15:1111.
Acta Materialia 57 (2009) 715–721 www.elsevier.com/locate/actamat
Effect of hydrogen on internal friction and elastic modulus in titanium alloys N.L. Arabajian a, V.I. Serdobintsev a, V.M. Tavkhelidze a,*, T.A. Peradze b, Yu.I. Stamateli b, K.M. Gorgadze b a
E. Andronikashvili Institute of Physics, 6 Tamarashvili St., 0177 Tbilisi, Georgia b Georgian Technical University, 77 Kostava St., 0175 Tbilisi, Georgia
Received 20 July 2008; received in revised form 4 October 2008; accepted 9 October 2008 Available online 10 November 2008
Abstract The effect of hydrogen on the variation with temperature of internal friction (QI) and elastic modulus (E) of a number of Ti-based alloys has been studied in the Hz and kHz frequency ranges. A relaxation peak of internal friction with a high degree of relaxation (QImax 101) and with a DE effect is observed in all hydrogen-doped samples at T 600 K at 1 kHz, and at T 500 K at 1 Hz. Such a peak is not present in samples without hydrogen. The activation energy W and the frequency factor v0 of the observed relaxation are determined to be W 1.55 eV, v0 1017 s1. It is shown that the observed effects are connected with the mechanism of grain boundary relaxation, as the introduction of hydrogen into titanium alloys leads to the formation of fine-grained structures. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Titanium alloys; Internal friction; Hydrogen; Grain boundaries relaxation
1. Introduction Titanium alloys with b-isomorphic elements (Nb, V, Ta, Mo, Zr) are structural materials that exhibit a shape memory effect (SME) and are used in different spheres of engineering. In contrast to the well-known and widely studied intermetallic alloys Ti–Ni, the materials investigated in this work are substitutional solid solutions in which the doping elements are disordered. In these alloys, the SME is based on thermoelastic phase transformation of body-centered cubic b-austenite into orthorhombic a-martensite [1]. At present, the following effects, which are connected with the introduction of hydrogen into titanium alloys, are known: a) decrease in the temperature of the martensite transition [2];
b) changes in the energy of stacking faults and in the energy of interphase and intergrain boundaries [3]; c) hydrogen-phase cold hardening [4]. Measurements of internal friction and elasticity modulus have shown that hydrogen exerts a significant influence on the elastic characteristics and damping ability of titanium alloys [2,5–7]. Although these effects connected with hydrogen are significant, their structural mechanisms are not yet quite clear. Thus, further investigation of the properties of hydrogenous Ti alloys is of a special interest. The present work deals with the investigation of the effect of hydrogen on the characteristics of the martensite transformation, elastic properties and damping ability of Ti alloys with Nb, V, Ta, Mo, Zr, over a wide range of temperatures and hydrogen concentrations. 2. Materials and methods
*
Corresponding author. E-mail addresses: [email protected], [email protected] (V.M. Tavkhe lidze).
Measurements were carried out on the alloys: Ti– 50.1Ta, Ti–59.8Ta, Ti–33Nb–7Zr, Ti–5.1Ta–4.9Mo–4.9 V,
1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.10.013
716
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
Ti–5.1Ta–4.9Mo–4.9 V–2.0Zr (all wt.%), obtained by multiple electron-beam remelting with subsequent hot rolling. The measurements of internal friction (QI) and elastic modulus (E) were carried out by the method of excitation of bending vibrations of the sample in the kHz frequency interval and torsional vibrations on a Coulomb-type pendulum in the frequency range 1–4 Hz. The internal friction was determined by the damping decrement of natural vibrations of the sample by means of an electronic discriminator registering the number of vibrations at half the vibration amplitude, and the elastic modulus was determined by the frequency f of resonance vibrations (according to the relation E / f 2 . The samples were thin plates 0.5 1 20 mm3. The initial phase state of the alloys was obtained by hardening in water after isothermal exposure in the temperature range (T > 1100 K) at which high-temperature b-austenite exists. Depending on the type of doping elements and their concentrations, the initial state of the alloy corresponded to (b00 +a00 )- or a00 -phase [8]. The characteristic temperatures of direct (Ms, start; Mf, finish) and inverse (As, start; Af, finish) martensite transformations were determined by differential thermal analysis (DTA) and from the curves of deformation – temperature dependence for torsional deformation under the action of permanent loading in the process of cooling the samples from a certain temperature above the martensite transformation and a subsequent heating. This method is applicable for determining the temperatures of the start and the finish of the martensite transformation owing to the SME. Introduction of hydrogen into samples was carried out in a special high-vacuum chamber at a temperature of 800 °C. After the chamber was filled with gaseous hydrogen, the change in hydrogen pressure was traced. Further, when the process of saturation of the sample with hydrogen was ended, the furnace in which the chamber with the sample was placed was rapidly removed, and the sample was cooled to room temperature in 10 s. The hydrogen concentration in the sample was determined by the pressure drop in the chamber during the hydrogenation process. 3. Experimental results 3.1. Ti–59.8 wt.% Ta alloy Fig. 1 shows the temperature spectra of internal friction and elastic modulus (in the units of sample frequency squared) of Ti–59.8 wt.% Ta alloy measured on a torsion pendulum with heating at the rate of 8 K min1. The form of the QI and f 2 curves is typical for phase transitions of the martensite type. As known [9], the internal friction at a martensite transformation is described by the relation:
Fig. 1. Temperature dependencies of (a) f2 and (b) QI of Ti–59.8 wt.% Ta alloy measured on a torsion pendulum during heating.
QI ðT Þ ¼
1 nE dT du ; 2 ra x dt dT
ð1Þ
where n is the shear deformation of the lattice at its structural rearrangement, E is the elasticity modulus, ra is the alternating voltage amplitude, T is the temperature, t is the time, x is the vibration frequency and u is the volume fraction of martensite phase. The measurements in the kHz frequency range give similar temperature dependencies for QI and f 2 at the excitation of the flexural mode. Fig. 2 (curve 1) shows the dependence QI (T) of the same sample measured at a frequency of 1 kHz. The sample heating rate was again 8 K min1. Comparison of Figs. 1 and 2 shows that the amplitude of the QI peak at the frequency of 4 Hz is about an order of magnitude higher than at the frequency of 1 kHz, though it does not fit Eq. (1). To check the dependence of peak height QI on the heating rate dT/dt, the experiment was carried out with a sharp change of T (curve 2 in Fig. 2). At TI = 480 K, the sample heater was switched off near the top of the peak, causing a sharp decrease in the heating rate, which drops to zero. As Fig. 2 shows, after stopping the heating, the increase in the internal friction sharply slows down and then decreases almost to the background value (the time of decrease is 15 min). The repeated heating at T2 = 500 K shows that QI begins to increase rapidly and practically joins curve 1 in Fig. 2. This observation proves the fact that the peak registered at the temperature of 500 K is indeed connected with the mar-
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
717
Fig. 2. Temperature dependencies of QI of Ti–59.8 wt.% Ta alloy measured in a flexural mode: curve 1, heating rate 8 K min1; curve 2, measurements in which heating was stopped at T1 = 480 K and resumed at T2 = 500 K.
tensite–austenite phase transition. The measurements by DTA give the following temperatures for the start and finish of the transition: As = 450 K, Af = 520 K. 3.2. Ti–59.8 wt.% Ta alloy doped with hydrogen After investigation of the initial samples of Ti–59.8 wt.% Ta alloy, hydrogenation was carried out according to the method described in Section 2. Three samples were prepared with hydrogen concentrations 0.07, 0.10 and 0.17 wt.% H. Fig. 3 shows the temperature dependencies of QI and 2 f of the Ti–59.8 wt.% Ta + 0.07 wt.% H sample measured during the process of heating at the excitation of the flexural mode. (It should be noted that the specific character of the measurement of QI using an electron discriminator does not unfortunately allow measurements of the values of QI 0.1 with an error <25%.) The shape of the curves points to the existence of a relaxation process with an unusually high degree of relaxation. The absence of the peak connected with martensite–austenite phase transition is also noteworthy, although the DTA measurements and the estimation of transition temperature by the method of sample bending show that the phase transition takes place in the temperature range 395–495 K. To estimate the activation energy of the relaxation process, the measurements of temperature dependence of QI and f 2 of the same sample were carried out on the torsion pendulum (Fig. 4). As seen from this figure, the temperature of the relaxation maximum almost coincides with the temperature of phase transition, and therefore it was checked whether the observed peak is caused only by relaxation or by the superposition of two processes: relaxation and phase transition. For this purpose, at the peak maximum, the temperature was stabilized and kept for 1 h; however, no changes in the level of vibration energy absorption were observed during this period.
Fig. 3. Temperature dependencies of (a) f2 and (b) QI of Ti–59.8 wt.% Ta + 0.07 wt.% H sample measured in a flexural mode in process of heating (As = 395 K, Af = 495 K).
Fig. 4. Temperature dependence of QI of Ti–59.8 wt.% Ta + 0.07 wt.% H sample measured on a torsion pendulum (j heating, h cooling).
In addition, a comparison of the QI (T) curves recorded during heating and cooling of the sample shows the absence of hysteresis in phase transitions of the first kind, which also indicates the existence of only a relaxation process. The estimation of activation energy of the process according to the temperature shift of the relaxation maximum in accordance with the Arrhenius equation gives the following values of activation energy W and frequency factor v0: W = 1.56 ± 0.02 eV, v0 = 1017 ± 0.5 s1. The measurements of QI (T) of Ti–Ta samples with 0.1 and 0.17 wt.% H show that the temperature position of the
718
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
observed relaxation peak does not depend on hydrogen concentration, and the degree of relaxation changes insignificantly (QImax 0.1 for 0.1 and 0.17 wt.% H concentration). 3.3. Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V alloy Fig. 5 shows QI(T) and f 2(T) dependencies of Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V alloy. The trend of the curves shows that in the temperature range 290– 770 K, any anomalies that can be connected either with the phase transition or with the relaxation are not observed, although the measurements of heat capacity (DTA) and SME show the martensite–austenite phase transition in the range of 670–750 K. One can assume that the peak of QI corresponding to the martensite–austenite transition is masked by the rapidly increasing background of internal friction. 3.4. Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H alloy Fig. 6 shows QI(T) and f 2(T) dependencies of Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H alloy measured at 1 kHz frequency. As is seen from the figure, the introduction of hydrogen into the sample changes the essential form of temperature spectra of QI and f 2.
Fig. 6. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H sample measured in a flexural mode in process of heating (As = 330 K, Af = 470 K).
As in the case of Ti–Ta + H system, in the range of T 610 K there appears a relaxation peak with a very high degree of relaxation (QImax 0.1) with a corresponding anomaly of the elastic modulus. The measurements at a low frequency (Fig. 7) allow us to calculate the activation energy of the peak W and the frequency factor v0 by the temperature shift of QImax. The obtained data W = 1.53 ± 0.02 eV and v0 = 1016±0.5 are in good agreement with the corresponding data for Ti–Ta + H system, which means that the nature of the observed peaks in both systems can be assumed to be identical. Measurements of the start (As) and finish (Af) temperatures of the martensite transition give values of As = 330 K and Af = 470 K. Thus, the introduction of hydrogen into the sample decreased the temperature of the martensite–austenite transition by about 200 K. This influence of hydrogen on the temperature of the martensite transformation was found earlier [2] in a study of industrial Ti-based alloys, and is explained by a stabilizing action of hydrogen on the high-temperature b-phase of the alloy. As in the case of the Ti–Ta + H system, the presence of the phase transition is not observed on the curve of QI(T) dependence. 3.5. Ti alloys of different compositions doped with hydrogen Fig. 5. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V sample measured in torsional vibration mode during heating (As = 670 K, Af = 750 K).
Apart from the measurements described above in detail, the measurements of the following alloys were also carried out: Ti–5.1 wt.% Ta–4.9 wt.% Mo–4.9 wt.% V–2.0 wt.%
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
Fig. 7. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V + 0.43 wt.% H sample measured in torsional vibration mode.
719
placed in a furnace heated to 800 °C to provide a rapid heating of the sample. The observation of the dynamics of pressure in the chamber has shown that within 10 min the sample is completely dehydrogenated. After reaching the ultimate vacuum (104 Pa), the furnace was quickly removed from the chamber, and the sample was cooled down to room temperature in 10 s. Fig. 8 shows the temperature spectra of QI (curve 1) and f 2 of the sample after degassing recorded at the flexural mode during the heating process. From curve 1 it is clear that the relaxation peak observed in the alloy with hydrogen is present again in the spectrum; however, its width is somewhat larger, and the temperature position of the maximum is shifted to lower temperatures. In addition, the behavior of the elastic modulus is not typical of a simple relaxation process, indicating, in our opinion, the reconstruction of internal structure of the sample. After heating up to 800 K during the measurement process and subsequent slow cooling, the measurements of QI (T) and f 2(T) were carried out again. The results are given in Fig. 8 (curve 2). The temperature position of the QI(T) peak, its amplitude and shape show a significant difference from the spectrum in curve 1. Checking the origin of the QI peak by stopping the heating near the temperature of the maximum has shown that the peak is caused by the martensite–austenite phase transition (Fig. 8, curve 3).
Zr + 0.3 wt.% H; Ti–33 wt.% Nb–7 wt.% Zr + 0.16 wt.% H; Ti–50.1 wt.% Ta + 0.21 wt.% H. The measurements have shown that after the introduction of hydrogen, in all these systems there appears a relaxation peak of internal friction of a high magnitude with the maximum at T 610 K at 1 kHz and at T 510 K at 1 Hz; hence it can be assumed that the mechanism of relaxation is the same for all systems studied here. 4. Discussion It is absolutely clear that the processes connected with the diffusion of interstitial hydrogen cannot be characterized by such high values of activation energy obtained in our experiments. For example, according to data [7] obtained for Ti– Ni–Cu alloy, the activation energy of relaxation peak connected with the motion of hydrogen is 0.6 eV. The same can be said about the relaxation processes of Snoek– Ko¨ster-type, which are connected with the ordering of hydrogen atoms considered as elastic dipoles in dislocation field, as well as about the processes of reorientation of substitutional atom–interstitial atom pairs [10]. To check the reversibility of the observed phenomena connected with the introduction of hydrogen into alloy, the studied Ti–59.8 wt.% Ta + 0.07 wt.% H sample was placed into a vacuum chamber for dehydrogenization. After reaching a vacuum of 104 Pa, the chamber was
Fig. 8. Temperature dependencies of (a) f2 and (b) QI of Ti–59.8 wt.% Ta sample after the dehydrogenation measured in a flexural mode. Curve 1, the first measurement during heating to T = 800 K; curve 2, the second measurement after heating and subsequent slow cooling; curve 3, measurement when the heating was stopped at T 500 K.
720
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
DTA and SME measurements confirm our statement. Just the same results are obtained by measuring the samples with 0.10 and 0.17 wt.% H after dehydrogenation. Hence, it can be assumed that the relaxation appearing after the introduction of hydrogen is connected with it indirectly; it is the result of the change of microstructure of the sample after hydrogenation. To explain such unusual behavior of the alloy, it is necessary to use the investigations carried out in Ref. [3] on Ti alloys with Al, which are similar in their properties to the alloys studied in the present work. In Ref. [3] it was observed that the introduction of hydrogen into samples by a method similar to ours leads to the appearance of a fine-grained structure. It is therefore logical to connect the effect discovered by us to grain boundary relaxation [10]. Although the existing mathematical models cannot describe exactly the grain boundary peaks of internal friction observed in many metals and alloys, the available experimental data [11–14] testify that in most cases the activation energy is close to 1.5 eV, and the degree of relaxation is of the same order (QImax 0.1) as in our experiments. Furthermore, according to the theory, the grain boundary peak is observed only when the grain size is less than the thickness of the sample (in our case 0.5 mm). It is also known that high-temperature annealing leads to an increase in the grain size (crystallites), and therefore the disappearance of the peak at T 600 K after heating up to T 800 K can be explained by the increase in the grain size up to >0.5 mm. The assumption of the formation of a fine-grained structure in the process of hydrogenation also explains the absence of the Q1(T) peak corresponding to martensite phase transition. Indeed, according to Eq. (1), the internal friction is proportional to the macroscopic shear strain n of the sample. On the other hand, as a result of hot rolling, the initial samples have a pronounced texture, owing to which the formed martensite plates (lamellas) are oriented predominantly in one definite direction. Therefore, the macroscopic shear strain makes a contribution to damping. As for the formation of a fine-grained structure as a result of hydrogenation, this causes isotropic arrangement of crystallites containing the martensite phase. Hence, the macroscopic shear strain does not arise as a result of martensite transition, and therefore the transition is not observed in Q1(T) spectra. After dehydrogenation and high-temperature annealing, the anisotropy in the orientation of martensite plates again becomes significant and leads to the appearance of the Q1 peak in the temperature range of martensite transformation. The dehydrogenation of Ti–5.1 wt.% Ta–4.9 wt.% Mo– 4.9 wt.% V + 0.43 wt.% H sample gives a result slightly different from that described above. Namely, after dehydrogenation at 800 °C and a rapid cooling down to room temperature, Q1(T) peak (Fig. 9) corresponding to grain-boundary relaxation decreases manifold and is almost indiscernible even at the first heating. This can be explained by a different kinetics of the increase of the grain size of Ti – 59.8 wt.% Ta and Ti–5.1 wt.% Ta–4.9 wt.%
Fig. 9. Temperature dependencies of (a) f2 and (b) QI of Ti–5.1 wt.% Ta– 4.9 wt.% Mo–4.9 wt.% V sample measured in flexural mode immediately after dehydrogenation.
Mo–4.9 wt.% V systems strongly differing from each other in the degree of doping by b-isomorphic elements. Thus, it can be stated that the hydrogen introduced into the studied alloys destroys the textured structure of martensite lamellae, leading to the appearance of a fine-grained structure with a high concentration of grain boundaries. Therefore, after introduction of hydrogen into the sample, the internal friction peak connected with the martensite transformation disappears, and a new relaxation peak appears due to the motion of grain boundaries. After the hydrogen is removed from the sample, the relaxation peak is still observed until the subsequent annealing restores the large-block structure of martensite lamellae. 5. Conclusions 1. It is shown that the introduction of hydrogen into the investigated Ti alloys causes the appearance of a Debye peak of internal friction with a high relaxation degree (Q1max 101), and the peak connected with the martensite phase transformation disappears. 2. The parameters of the discovered relaxation process, activation energy W and frequency factor v0, indicate that the relaxation cannot be connected with the process of hydrogen atoms jumping between interstitial positions in the sublattice. 3. It is shown that the discovered relaxation is indirectly connected with hydrogenation of samples and can be
N.L. Arabajian et al. / Acta Materialia 57 (2009) 715–721
explained by the processes of transformation of the crystal lattice from a large-block-textured state into an isotropic fine-grained state. 4. The analysis of the obtained data leads to the conclusion that the observed QI peak is caused by the mechanism of grain boundary relaxation, which is the result of viscous slipping of neighboring grains with respect to each other. 5. The impossibility of observing the martensite peak of internal friction in hydrogenated samples is also connected with the formation of a fine-grained structure and, as a consequence, with the absence of macroscopic shear strain in the process of martensite transformation. Acknowledgments The authors thank Dr. V. Melik-Shakhnazarov for his valuable comments and N. Goldbaum for her technical assistance in preparing the manuscript.
721
References [1] Zvicker U. Titan and titanlegierungen. Berlin, Heidelberg, New York: Springer-Verlag; 1974. [2] Ilyin A, Kollerov M, Golovin I. J Alloys Comp 1997;253:144. [3] Kolachev B, Nosov V. FMM 1984;57(2):286 [in Russian]. [4] Goltsov V. Mater Sci Eng 1981;49(2):109. [5] Mazzolai FM, Biskarini A, Campanella R, Collusi B, Mazzolai G, Rotini A, et al. Acta Mater 2003;51:573. [6] Biskarini A, Collusi B, Mazzolai G, Tuissi A, Mazzolai FM. J Alloys Comp 2003;355:52. [7] Fan G, Zhou Y, Otsuka K, Ren X, Nakamura K, Ohba T, et al. Acta Mater 2006;54:5221. [8] Lyasotskaya V, Lyasotskyi I, Ravdonikas I, Belov S. In: Proceedings of institutes of higher education, vol. 3, 1985. p. 82. [9] Malygin GA. Physics-Uspekhi 2001;44(2):173. [10] Novick AS, Berry BS. Anelastic relaxation in crystalline solids. New York: Academic Press; 1972. [11] Ke T. Phys Rev 1947;71:533. [12] Weinig S, Mashlin E. Trans Met Soc AIME 1957;209:32. [13] Cordea J, Spretnak J. Trans Met Soc AIME 1966;236:1685. [14] Williams T, Leak G. Acta Met 1967;15:1111.