Tensile properties and high temperature creep behavior of microalloyed Ti–Ti3Al–Nb alloys by directional solidification

Materials Science and Engineering A 527 (2010) 4484–4496

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Tensile properties and high temperature creep behavior of microalloyed Ti–Ti3 Al–Nb alloys by directional solidification Qingdong Luan, Qiuqi Duan, Xiaoguang Wang, Jing Liu, Liangming Peng ∗ CAS Key Laboratory for Mechanical Behavior and Design of Materials, Department of Modern Mechanics, School of Engineering Science, University of Science and Technology of China, Hefei 230026, PR China

a r t i c l e

i n f o

Article history: Received 16 September 2009 Received in revised form 23 March 2010 Accepted 26 March 2010

Keywords: Directional solidification Microalloyed Ti–Ti3 Al–Nb alloys Tensile properties Creep deformation Dynamic recrystallization

a b s t r a c t The microstructures, tensile properties and creep behavior of the Cr- and Mo-microalloyed Ti–Ti3 Al–Nb alloys processed by directional solidification were investigated. The experimental results indicated that the strength, ductility and creep resistance of ternary Ti–Al–Nb alloy were remarkably enhanced by additions of small amounts of Cr and Mo. The stress exponent for creep showed a transition and varied from a lower value of 3.5–5.3 (LSR-region I) to a higher value of 8.3–12.6 (HSR-region II) with increasing stress at 600 and 650 ◦ C whereas a single stress exponent value of 5.3–6.0 was obtained at 700 ◦ C. Conversely, the true activation energy for creep was calculated to be 303–324 kJ/mol and fell into or was quite close to that for the self-diffusion of Ti in ␣2 -Ti3 Al (288–312 kJ/mol). TEM examinations revealed that ordinary dislocations in the lath dominated the deformed microstructures. However, the initial microstructure of the alloys was unstable during long-term creep exposure and dynamic recrystallization occurred at high stress and moderately high temperature. The creep deformation in region I was dislocation-controlled whereas the abnormally high stress exponent in region II was associated with the effects of recrystallization resulting in a reduction in the overall grain size of the initial structure. The fracture modes and total strain were dependent on the composition, test temperature, and stress level. The accelerating strain rate in the extended tertiary stage was attributed to the microstructural instabilities or the nucleation, coalescence and merge of voids or microcracks. © 2010 Elsevier B.V. All rights reserved.

1. Introduction There has been considerable recent interest in the development of advanced titanium aluminide-based alloys with a compositiondependent structure of single-phase Al-rich ␣2 -Ti3 Al, Ti-rich ␥-TiAl or alternating ␥-TiAl and ␣2 -Ti3 Al lamellae. These alloys are currently considered as potential materials for high temperature applications owing to their high strength/density ratio and excellent creep resistance. For instance, the Ti3 Al-based alloys can retain reasonable strength levels up to 800 ◦ C. Nevertheless, they suffer from poor ductility and fracture toughness. Consequently, recent development of titanium aluminide alloys has been directed to the combinations of room temperature ductility and elevated temperature creep strength mainly in view of two approaches. The first one is alloying approach via the additions of ␤-phase stabilizing elements such as Nb, Mo, V, Re and W to the binary Ti–Al alloys, resulting in Ti–Al–X (X denotes the aforementioned ␤ stabilizers) systems with improved properties [1–5]. Another one is focused

∗ Corresponding author. Tel.: +86 551 360 6964; fax: +86 551 360 6459. E-mail address: [email protected] (L. Peng). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.03.096

on the control of the orientation of the lamellar microstructure in Ti–Al alloys by the directional solidification (DS) process [6–10]. In particular, when the lamellar boundary orientation was aligned parallel to the tensile direction, the Ti–Al alloys could possess a reasonable combination among creep resistance, yield strength and tensile elongation [1,2,11,12]. The systematic understanding of creep behavior in these materials is of utmost importance in view of their potential for elevated temperature applications. So far, extensive investigations have been directed to binary or multi-microalloyed Ti–Al alloys with Al concentration in the range of 25–53 at.% [3,5,13–28]. Table 1 summarizes the reported values of stress exponents, activation energies and possible operative mechanisms for creep deformation in Ti–Al–M alloys with different microstructures under different test conditions. It is evident that the measured values of the stress exponent, n are in the wide range of 2.3–19 and the apparent activation energies for creep, Qa range from 0.65 to 2.17 times of the value for Ti self-diffusion in single ␥-TiAl. Correspondingly, several mechanisms operative in the creep deformation of these Ti–Al–M alloys have been proposed in the literature including dislocation activity (glide, climb and jog) [3,13–18], grain-boundary sliding [19–21], mechanical twinning [22–24], and dynamic recrystallization [25–27], which are quite sensitive to the constituting phases,

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microstructural morphologies, and test conditions. In particular, several creep studies indicated a transition in the stress exponent and activation energy, and in turn, a transition in mechanism for creep deformation in titanium aluminides [16,26]. It is obvious that previous studies have been mainly concentrated on the creep behavior of single ␥-TiAl alloys, duplex or fully lamellar ␥ + ␣2 two-phase alloys. Unfortunately, quite few investigations have been undertaken up to date on the creep behavior in single ␣2 -Ti3 Al alloys [29–32]. Furthermore, almost no reports have been available up to date with respect to the high temperature creep deformation in the ternary Ti–Al–Nb alloys with a moderate Al concentration (10–20 at.%), where some of the remaining Ti is employed as the tough phase to enhance the ductility of Ti–Al–Nb alloys resulting in a composite microstructure of Ti–Ti3 Al–Nb. On

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the other hand, it has been indicated that the addition of Cr is mostly effective in enhancing the ductility whereas the Mo addition promotes the ␤ solidification process and simultaneously provides some solid solution strengthening in TiAl alloys with Al content ranging from 42 to 49 at.% [1,33–35]. In this study, the original concept of developing the Ti–16Al alloys with a composite microstructure is derived from the improvement in the ductility and creep resistance of Ti3 Al alloys by remaining Ti phase as well as through adding large amounts of Nb and microalloyed elements, Cr and Mo. The directional solidification process was employed to fabricate these alloys. Their tensile and creep deformation were investigated including experimental data and microstructral changes during creep with an attempt to elucidate the fundamental mechanisms controlling the

Table 1 Reported values for n, Qa and possible mechanisms operative in creep of Ti–Al–M alloys (in atomic composition) with different microstructures. Alloy and microstructure studied

Stress exponent, n

Activation energy, Qa (kJ/mol)

Test conditions

Possible mechanism

Reference

Ti–48Al: equiaxed near ␥

5–6

–

760 ◦ C; 100–300 MPa

Viswanathan et al. [13–15]

Ti–47Al–4(W, Nb, B): nearly-lamellar

5–7

–

760 ◦ C; 100–500 MPa

Ti–46Al–2W–0.5Si: lamellar (␣2 + ␥) + feathery (␣2 + ␥) + precipitates of Ti5 Si3 and Ti-based solid solution Ti–40Al–10Nb: Widmanstätten lath Ti–48Al: duplex-TiAl

7.3

405

700–800 ◦ C; 200–390 MPa

Climb of 1/2[1 1 0]-type jogged-screw dislocations. Dislocation climb at low stresses; Jog-dragged dislocation glide at high stresses. Non-conservative motion of dislocations.

3

359 –

Ti–48Al–1.5Cr: nearly-lamellar

2.4–3.5 (LSR); 8.6–14.3 (HSR)

292 (LSR); 630 (HSR)

700–900 ◦ C; 70–400 MPa

Ti–46Al–2W–0.5Si: duplex-TiAl-equiaxedprimary ␥ + lamellar (␥ + ␣2 ) colonies Ti–25Al: single equiaxed ␣2 -Ti3 Al

5 (LSR) 10 (HSR)

290 (LSR)

700 and 750 ◦ C; 225–370 MPa

Dislocation glide and climb. A threshold stress resulting from the mechanical twinning at interfaces. Dislocation glide-controlled creep at lower stresses and temperatures; Dynamic recrystallization controlled creep at higher stresses. Dislocation climb in LSR.

Zhan et al. [18]

19

750 ◦ C; 160–240 MPa 700 ◦ C; 250–430 MPa

2.3–2.5 ( < 138 MPa, T ≥ 750 ◦ C); 4.3–5 ( < 138 MPa)

206

650–800 ◦ C; 70–300 MPa

Dislocation climb at high temperatures and stresses; Grain-boundary sliding at low temperatures and stresses.

Mendiratta and Lipsitt [29]

5.5 (T < 650 ◦ C); 6.5 ( > 172.5 MPa, T ≥ 650 ◦ C); 2.5 ( < 172.5 MPa, T ≥ 650 ◦ C) 4.5 (LSR); 5.8 (ISR); 4.9 (HSR)

190 285

600–800 ◦ C; 69–312 MPa

370 380 340

777–977 ◦ C; 70–500 MPa

Dislocation glide and/or the dislocation recovery at LSR; Deformation twins at HSR.

Nagae et al. [30]

Ti–26Al–5Nb: single equiaxed ␣2 -Ti3 Al with Nb solid solution

Ti–34Al: single equiaxed ␣2 -Ti3 Al

LSR, low stress region; ISR, intermediate stress region; HSR, high stress region.

Zhang and Deevi [16]

Lapin [17]

Morris and Leboeuf [22]

Allameh et al. [26]

Spigarelli et al. [28]

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Fig. 1. Schematic drawing of the directional solidification unit.

intermediate temperature creep deformation in Cr and Mo-doped Ti–Ti3 Al–Nb alloy. 2. Experimental procedures 2.1. Material processing and sample preparation Ingots of Ti–16Al–12Nb and Ti–16Al–12Nb–3Cr–1Mo (in at.%) alloys were prepared using pure-Ti (99.99 mass%), Al (99.99 mass%), Nb (99.99 mass%), Cr (99.9 mass%) and Mo (99.9 mass%) at a highpurity argon atmosphere in a vacuum electromagnetic induction furnace with a water-cooled copper crucible. The melting was repeated more than 5 times to ensure homogeneity. Then, the master ingots were re-melted in a vacuum induction furnace using a CaO crucible. The DS ingots were grown in a laboratory scale vacuum furnace using Ti–16Al–12Nb alloy as a seed material at a growth rate of 60 mm/h. The directional solidification unit was schematically illustrated in Fig. 1 and the furnace temperature was heated to the range of 1760–1820 ◦ C maintaining the vacuum at 1.0 Pa level. Flat dog-boned specimens, 25 mm in gauge length and 5.5 mm × 2 mm in cross-section, were electric-discharge machined with the tensile loading axis parallel to the growth direction. All the specimens for tensile and creep tests were mechanically ground to 1500 grit sand paper and then electropolished with a voltage of 6 V in a solution of 40 ml H3 PO4 and 10 ml HF to eliminate the deleterious influence of surface defects on mechanical properties.

at temperatures of 600, 650 and 700 ◦ C under applied stresses of 100–440 MPa on the same testing machine. The temperature of a specimen was measured with a temperature accuracy of ±1 ◦ C by three thermocouples closely attached to the upper, middle and lower section of the specimen, respectively. Both the tensile and the creep strains were continuously measured using a Linear Variable Differential Transducer (LVDT) extensometer having a strain resolution of ±0.1 ␮m. The acquisition of time-elongation data was accomplished by a computer and data processing was conducted through a computer program. Creep rates were determined at 6–8 different stress levels for each temperature. 2.3. Microstructural characterization and fracture morphology The longitudinal sections of specimens were polished and etched in a solution of 5 ml HF, 15 ml HNO3 , and 50 ml H2 O. The microstructures before and after creep tests and fracture surfaces were examined using scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS). The creep samples were quickly cooled in air after the creep tests were interrupted or finished. Thin foils for transmission electron microscopy (TEM) observation were cut from the crept specimen gauge sections prepared by discs sectioned parallel to the stress axis. The foils were mechanically polished from 0.4 mm to approximately 0.1 mm and finally thinned by twin-jet polishing at 15 V and −30 ◦ C in a solution consisting of 195 ml methanol, 150 ml n-butyl alcohol and 15 ml perchloric acid. TEM observations were conducted on a JEOL JEM3010 microscope operated at 200 kV.

2.2. Tensile and creep tests 3. Results The conventional tensile tests at room and high temperatures (650 and 750 ◦ C) were carried out at a constant strain rate of 1.0 × 10−4 s−1 in air on a CSS-3905 multi-functional electronic relaxation creep testing machine equipped with a three zone furnace. Constant-load tensile creep tests were performed

3.1. Microstructures before creep Fig. 2 shows SEM second-electron images of longitudinal sections for the DS Ti–16Al–12Nb (Fig. 2a) and

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Fig. 2. Second-electron images of longitudinal section and EDS spectra for (a) Ti–16Al–12Nb and (b) Ti–16Al–12Nb–3Cr–1Mo alloys to indicate the chemical composition and constitute phases.

Ti–16Al–12Nb–3Cr–1Mo (Fig. 2b) alloys. Both the two alloys consisted of Widmanstätten laths in a Ti-rich matrix. It is evident that the addition of Cr and Mo reduced the inter-lamellar spacing and almost the lath orientation of the two alloys was not parallel to the withdrawal direction. Instead, most of them tended to be inclined with a respective angle of approximately 45◦ and 30◦ to the withdrawal direction. This indicates that the Mo and Cr additions would promote the lamellar microstructure growth since the slow diffusion rate of Mo may be beneficial for improving the stability of the lamellar microstructure against spheroidization during solidification [10]. The EDS analysis and XRD pattern results show that the remaining Ti phase in the alloys exhibited three morphologies: lath, block, and equiaxed. It should be noted that no formation of Ti2 AlNb (O phase) was found in the present alloys, which is different from the Ti–(18–30)Al–(12.5–30)Nb (in at.%) alloys where the Ti2 AlNb serves as the major phase [36–38]. Instead, a small part of Nb was dissolved into the Ti matrix and the rest appeared in the Nb-rich compounds, Nb2 Al and Cr2 Nb whereas most of the added Mo reacted with Al to yield small amount of aluminide compound, Al5 Mo throughout the pre-solidified Ti matrices. 3.2. Tensile properties Fig. 3 shows the true stress vs. true strain curves at room and high temperatures for DS Ti–16Al–12Nb and Ti–16Al–12Nb–3Cr–1Mo DS alloys and their maximum tensile stress and elongation to failure are summarized in Table 2. It can be found that both the two alloys exhibited little ambient Table 2 Maximum tensile stress and elongation to failure for the two alloys at room and high temperatures. Alloy

Ti–16Al–12Nb Ti–16Al–12Nb–3Cr–1Mo

Maximum tensile stress (MPa)

Elongation to failure (%)

RT

650 ◦ C

750 ◦ C

RT

650 ◦ C

750 ◦ C

509 559

378 454

260 258

0.4 0.6

0.8 3.9

5.0 19.6

Fig. 3. Tensile true stress–true strain curves at different temperatures for (a) Ti–16Al–12Nb and (b) Ti–16Al–12Nb–3Cr–1Mo alloys.

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temperature ductility, and splitting type fracture immediately occurred after the stress reached the maximum without obvious yielding process. An insignificant yielding process was observed in Ti–16Al–12Nb alloy at 650 ◦ C. While for Ti–16Al–12Nb–3Cr–1Mo alloy tested at the same temperature, the stress increased gradually with increasing strain after yielding and finally a plateau appeared, suggesting a balance between the working hardening and the recovery softening processes. During tensile tests at 750 ◦ C, they experienced remarkable plastic deformation. In particular, obvious necking took place for Ti–16Al–12Nb–3Cr–1Mo alloy prior to final fracture. In general, the maximum tensile true stress decreased and the elongation increased with increasing temperature. The Ti–16Al–12Nb–3Cr–1Mo alloy exhibited higher maximum tensile stress than Ti–16Al–12Nb alloy at room temperature and 650 ◦ C. Furthermore, the maximum stress increased with the content of Nb for the same alloy system. For instance, the maximum stress increased from 509 MPa for Ti–16Al–12Nb to 559 MPa for Ti–16Al–12Nb–3Cr–1Mo alloy at room temperature. While at 650 ◦ C, the maximum stress took the values of 378 MPa vs. 454 MPa for the former and latter alloys, respectively. The strength increase might be partially associated with the fact that, as mentioned previously, the majority laths of Ti–16Al–12Nb–3Cr–1Mo tended to be inclined with a smaller angle to the growth (tensile) direction than Ti–16Al–12Nb alloy. Nevertheless, the difference in maximum stresses between them disappeared at 750 ◦ C. Similarly, the ductility enhancement was also achieved in Ti–16Al–12Nb–3Cr–1Mo alloy. The tensile elongation was enhanced from 0.4 to 0.6% and from 0.8 to 3.9% at room temperature and 650 ◦ C, respectively. Comparatively, a much more significant enhancement in the elongation at 750 ◦ C, i.e. 19.6% was achieved. The present results were consistent with those reported in the literatures, where the Cr addition improves the ductility and Mo addition contributes an improvement in the strength of Ti–Al alloys [1,33–35]. 3.3. Creep deformation Fig. 4 shows the typical curves of creep strain vs. time for the two alloys at 600, 650 and 700 ◦ C at different stress levels. As evident from Fig. 4a, the upward creep curves at 600 ◦ C consisted of three regions: primary, steady-state and tertiary stages. Nevertheless, the tertiary stage was quite short and creep fracture occurred upon onset of this process. In contrast, the downward creep curves at 650 ◦ C exhibited short primary creep that was directly followed by a long-term steady-state creep as well as an obvious tertiary creep stage. As the testing temperature increased to 700 ◦ C, almost no steady-state creep stage was observed. In this case, the creep rate decreased with strain or time in the initial portion of the curve, reached a minimum value after the primary stage. Then the creep deformation entered into an extended region where the creep rate continually increased with time. The overall amount of creep strain was observed to increase with temperature. Simultaneously, the figures clearly presented that the Cr- and Mo-containing Ti–Al–Nb alloys were more creep resistant and exhibited larger tertiary creep strain than that of ternary Ti–Al–Nb alloy. It has been well established that the relationship between the steady-state (minimum) creep rate, ε˙ in polycrystalline materials and applied stress,  can be described by the Dorn phenomenological constitutive equation [39]: ε˙ =

 p   n AD0 Gb b kT

d

G

 Q

exp −

RT

,

(1)

where A is a dimensionless constant, D0 is a frequency factor included in the term for diffusion coefficient, d is the grain size of polycrystalline materials, Q is the activation energy for creep deformation, p and n (=4–7) are the grain size and stress exponent, respectively. The remaining terms have been defined previously.

Fig. 4. Creep strain vs. time for the creep tests of the two alloys at different temperatures. (a) 600 ◦ C, (b) 650 ◦ C, and (c) 700 ◦ C.

By fitting the data to the above equation, the stress exponent, n can be determined. The variations of the steady-state (minimum) creep rate with the applied stress on double logarithmic scales were shown in Fig. 5, where the slope exhibited a stress-dependent transition and two distinct regions, I and II with different stress exponents were observed at 600 and 650 ◦ C for the two alloys. In region I, a stress exponent between 3.5 and 5.3 was observed for low stress levels, whereas stress exponent values between 8.3 and 12.6 were obtained in region II for higher applied stresses.

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Fig. 5. Variations of the steady-state (minimum) creep rate with the applied stress for (a) Ti–16Al–12Nb and (b) Ti–16Al–12Nb–3Cr–1Mo alloys. The stress exponent and correlation coefficient for the fits were indicated in the figures.

However, there was only one region in the curves at 700 ◦ C and the stress exponents (n) were 6.0 and 5.3 for Ti–16Al–12Nb and Ti–16Al–12Nb–3Cr–1Mo alloys, respectively. It is worthy noting that the correlation coefficients, r of all these linear fits were higher than 0.989. Based on Eq. (1), the true activation energy, Q for creep can be calculated according to the following equation:



Q =

R∂ ln ε˙ ∂(1/T )



+ nR 

T 2 ∂G . G ∂T

(2)

However, in view of experimental operation, the activation energy can be measured based on the first term of the above equation for a given stress by neglecting the second small term and plotting log ε˙ against the inverse of temperature, 1/T. In this way and as shown in Fig. 6, the activation energy values were calculated to be approximately 340 kJ/mol without considering the temperature dependence of the shear modulus (the second term in Eq. (2) usually takes a value of only 10–20 kJ/mol) for both regions I and II whereas only two data points were available in most cases. It can be found that the measured stress exponents can be compared with those reported in the literatures for various single ␥-, lamellar-, duplex-TiAl, and Ti3 Al alloys [13–28], where the values ranged from 2.3 at low stresses and to as high as 19 at high stresses in some of these alloys. It should be noted

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Fig. 6. Arrhenius plot to determine the activation energy for LSR (region I) and HSR (region II).

that the transition in the stress exponent for creep has been also frequently observed in some of the aforementioned alloys. For instance, the extruded nearly-lamellar Ti–48Al–1.5Cr alloys tested in a temperature range of 700–900 ◦ C under 70–400 MPa by Allameh et al. [26] exhibited a transition in the stress exponent from 2.4–3.5 in the low stress region to 8.6–14.3 in the high stress region. While for the powder-metallurgy (PM) Ti–47Al–2Cr–2Nb and Ti–47Al–2Cr–1Nb–1Ta alloys with fine lamellar structures, the stress exponent increased with increasing stress from 1.5 to 10 [16]. Such a transition also occurred in some Ti3 Al-based alloy such as Ti–25Al, Ti–26Al–10Nb [29] and Ti–34Al [30] alloys. On the other hand, similar to lamellar Ti–48Al and duplex Ti–48Al–2Mn–2Nb alloys with a stress exponent of 19 above 310 MPa at 700 ◦ C and lamellar Ti–48Al–1.5Cr with the highest value of n of 14.3, the present alloys also exhibited such a high n value between 8.3 and 12.6 at 600 and 650 ◦ C in the high stress region. Unlike the stress exponent, the activation energies for creep took almost identical values and exhibited no transition in the LSR and HSR. The present result was somewhat beyond expectation and inconsistent with those reported in the previous studies where the activation energy ranged from 190 to 630 kJ/mol and in some cases a transition existed [18]. It is documented that the true activation energies for self-diffusion of Ti and Al in ␣2 -Ti3 Al are 288–312 kJ/mol (QTi-␣2 ) and 395 kJ/mol (QAl-␣2 ) [40], respectively. In contrast, they take lower values of 290 [41] and 360 kJ/mol [42]

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˙ Fig. 7. Normalized creep rate, εkT exp(Q/RT )/G (Q = 340 kJ/mol and G(GPa) = 53.56–0.0180T [45]) as a function of the normalized applied stress, /G for (a) Ti–16Al–12Nb and (b) Ti–16Al–12Nb–3Cr–1Mo alloys. A transition in stress exponent was still obviously observed indicating two regions.

for self-diffusion of Ti (QTi-␥ ) and Al (QAl-␥ ) in single ␥-TiAl while the value for self-diffusion of pure ␣-Ti, Q␣-Ti is 242 kJ/mol [43]. Nevertheless, it should be pointed out that the activation energy for self-diffusion in Ti–Al-based and Ti alloys usually spans a wide range and the present value of 340 kJ/mol was in a well agreement with the majority of the literature findings, as summarized in Table 1. It is evident that the present value was slightly higher than QTi-␣2 prior to taking the temperature dependence of shear modulus into account. Moreover, this value for the Ti–Ti3 Al–Nb (Ti–16Al–12Nb system) alloys is quite close to that (332 kJ/mol) for creep in the Ti3 Al–Nb ally (Ti–25Al–10Nb in at.%) with a basketweave microstructure [44]. In order to eliminate the influence of temperature on creep rate and verify the strengthening effect of the microalloying ele˙ ments, the normalized strain rates εkT exp(Q/RT )/G were plotted against the shear modulus-compensated stress /G in Fig. 7, where Q = 340 kJ/mol, and the dependence of the shear modulus (G) on the test temperature T (in ◦ C) was taken as G (GPa) = 53.56–0.0180T for Ti–16Al–10Nb alloy [45]. As evident from this figure, two regions appeared again in all the curves and the correlation coefficient r for all the fits was higher than 0.97. All of the data at different temperatures for the individual alloy fell very close to a straight line indicating that it was appropriate to take Q = 340 kJ/mol in the

Fig. 8. Structural morphology after a long-time creep under low stresses for (a) Ti–16Al–12Nb (650 ◦ C/250 MPa/40 h) and (b) Ti–16Al–12Nb–3Cr–1Mo (600 ◦ C/210 MPa/140 h). Note that some vacant traces were left where the presolidified Ti blocks have autolyzed into the matrices.

temperature range of 600–700 ◦ C. Such a data processing yielded stress exponents, n of 6.3 and 5.2 in LSR, 8.9 and 11.5 in HSR for Ti–16Al–12Nb, and Ti–16Al–12Nb–3Cr–1Mo, respectively. It can be found that the creep resistance of the Ti–16Al–12Nb–3Cr–1Mo alloy was superior to that of the Ti–16Al–12Nb alloy since its creep rate was lower at a given (normalized) stress. Previous studies have also demonstrated that the additions of refractory elements such as Mo, W and Ta can improve the creep resistance of Ti–Al alloys owing to their solid solution strengthening and low diffusivity [43,46]. 3.4. Crept structures Fig. 8 illustrates the microstructures of longitudinal sections near the fracture surfaces for the different specimens subjected to long-term creep tests. Compared with the initial microstructures in Fig. 2, it was found that no cavities have nucleated in these regions. However, some pre-solidified Ti blocks have autolyzed into the matrices as indicated by the arrows where the vacant traces were left. In general, the amount of Ti blocks decreased implying that localized dynamic recovery and recrystallization has occurred during the long-term high temperature exposure under applied stresses. The dislocation structures in the alloys were examined to understand the creep mechanism using TEM after creep under different conditions. The results for the Ti–6Al–12Nb alloy are shown in Fig. 9. It is evident that the overall dislocation density in the sam-

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ple after crept at 600 ◦ C and 250 MPa (in LSR) was relatively low in the deformed section (Fig. 9a) and only a few of short straight dislocations were displayed inside the lath. The slightly curved and randomly distributed dislocation configurations at higher stress level were observed in Fig. 9b for a sample deformed at 600 ◦ C and 420 MPa (in HSR) although the dislocation density was still low. However, there was a significant increase in dislocation activ-

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ity for a sample deformed at 650 ◦ C under a respective stress of 190 and 320 MPa, where dislocation network trended to form inside the laths and at the colony boundaries (Fig. 9c and d). While after crept at 700 ◦ C and 160 MPa a large amount of free dislocation segments was observed as shown in Fig. 9e. Fig. 10 shows the deformation microstructures in the Ti–16Al–12Nb–3Cr–1Mo alloy subjected to creep at 600 ◦ C. Sim-

Fig. 9. TEM micrographs showing the deformed microstructures in the Ti–16Al–12Nb alloy. (a) Low density of short straight dislocation segments subjected to creep at 600 ◦ C and 250 MPa (LSR) for 40 h to a strain of 1.6%; (b) slightly curved and randomly distributed dislocation configurations after creep at 600 ◦ C and 420 MPa (HSR) for 1.6 h to a strain of 2.5%; (c and d) high density of dislocations with the trend of network formation located inside the laths and at the colony boundaries after creep at 650 ◦ C under 190 and 320 MPa, respectively; and (e) high density of free dislocation segments after deformed at 700 ◦ C and 160 MPa for 44 h to a strain of 18.6%.

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Fig. 10. TEM micrographs showing the deformed microstructures in the Ti–16Al–12Nb–3Cr–1Mo alloy subjected to creep at 600 ◦ C. (a) Low density of short parallel dislocation segments inside laths under 210 MPa (LSR) for 18 h to a strain of 1% without final failure and (b) high density curved dislocations under 380 MPa for 2 h to a strain of 1.5%.

ilarly, quite low density of short parallel dislocation segments was observed inside laths under 210 MPa (LSR) for 18 h to a strain of 1% without final failure (Fig. 10a). The density of curved dislocations increased with increasing applied stress and creep strain (HSR: 380 MPa for 2 h to a strain of 1.5%). The deformed microstructures in the Ti–16Al–12Nb–3Cr–1Mo alloy subjected to creep at 650 ◦ C were showed in Fig. 11. Again, few dislocations were visible in Fig. 11a under 280 MPa (LSR) for 4.2 h to a strain of 4.1%. Fig. 11b revealed strong evidence of dynamic recrystallization during the creep process at the same temperature and 310 MPa (in HSR). Under such a circumstance, polygonization of grains occurred generally resulting from the thermal-activation-aided climb of edge dislocations. Many sub-grains and low angle sub-boundaries were formed in this sample as indicated in Fig. 11c where arrays of dislocations similar to that generally observed in pile-ups. In general case, the sub-boundaries consisted of scores of short parallel dislocations belonging to a group or several groups in high temperature creep for alloys although the dislocation character was not analyzed in this study. The dislocation substructures in a lath in the Ti–16Al–12Nb–3Cr–1Mo specimen after creep to 20.9% strain at 700 ◦ C under 180 MPa were presented in Fig. 12. Moderately high density of tangled or bowed dislocations was observed in regions around the grain-boundary (Fig. 12a and b). In addition, short paral-

Fig. 11. TEM micrographs showing the deformed microstructures in the Ti–16Al–12Nb–3Cr–1Mo alloy subjected to creep at 650 ◦ C. (a) Low density dislocations after creep test at 650 ◦ C and 280 MPa (LSR) for 4.2 h to a strain of 4.1%; (b) dynamic recrystallization of laths and (c) short parallel dislocation pile-ups near grain boundaries after creep test at 650 ◦ C and 310 MPa (HSR) for 17.5 h to a strain of 9.7%.

lel dislocation pile-ups were also observed at the phase boundaries (Fig. 12c). On the whole, nearly no twinning and jogged dislocations were observed in the samples, which is different from those frequently reported in some ␣2 -Ti3 Al- and TiAl-based alloys deformed at high stresses [13–15,22,24,30,47].

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of fracture surfaces at 600 ◦ C under low stress was quite flat and transgranular cleavage fracture (Fig. 13a). Some secondary small cracks were observed on the fracture surface when the sample was applied to high stress (Fig. 13b). Conversely, the feature of fracture surface of the sample tested at 700 ◦ C and low stress was ductilelooking and creep cavities were formed during the long-term creep deformation as shown in Fig. 13c. Many large cracks were present on the fracture surface when high stress was applied to the sample (Fig. 13d). The Ti–16Al–12Nb–3Cr–1Mo alloy exhibited the similar fracture features (Fig. 14a–d). However, some tearing striations were frequently observed. In general, the fracture modes and total strain prior to final failure were dependent on the composition, test temperature, and stress level. The Cr- and Mo-microalloyed alloys exhibited high total strain, indicating that a significant plastic deformation even necking took place before final fracture. As evident from Fig. 4b and c, a considerable part of the creep life was consumed in the tertiary creep stage, which was mainly attributed to the fact that the high temperature and stress eventually led to microstructural degradation in the form of spherioidization or the recrystallization of lath and pre-solidified Ti blocks. This, in turn, enhanced the grain boundaries’ sliding, resulting in quite an early onset of the tertiary creep. During this stage, the strain rate was rapidly increased along with increasing accumulated strain, resulting in unstable deformation and leading to the damage of alloys. The accelerating strain rate was commonly connected with microstructural instabilities or the nucleation, coalescence and merge of voids or microcracks as have been shown in the above fractographs.

4. Discussion 4.1. Effects of Cr and Mo on microstructure and mechanical properties

Fig. 12. Dislocation structures in the crept Ti–16Al–12Nb–3Cr–1Mo alloy at 700 ◦ C and 180 MPa for 29 h to a strain of 20.9%. (a and b) Tangled or bowed dislocations and (c) short parallel dislocation pile-ups near phase boundaries.

3.5. Fractographs and fracture analysis Figs. 13 and 14 display selected fractographs of Ti–16Al–12Nb and Ti–16Al–12Nb–3Cr–1Mo alloys after the creep tests under different conditions, respectively. Fig. 13a and b is the respective fracture surfaces for Ti–16Al–12Nb alloy subjected to creep at 600 ◦ C under 250 and 460 MPa while Fig. 13c and d for 700 ◦ C under 160 and 320 MPa, respectively. It is evident that the characteristic

Based on the Ti–Al binary phase diagram as shown in Fig. 15, the solidification process of the Ti–16Al alloy follows the transformation sequence of L → ␤ → ␤ + ␣ → ␤ + ␣ + ␣2 → ␣ + ␣2 as indicated by the long short-dotted arrow. When the alloy was cooled to the two-phase region of ␤ + ␣, the ␣ phase precipitated from the ␤ matrix through the procedure of Widmanstätten precipitation. Subsequently, when the alloy was further cooled to the three phase region of ␤ + ␣ + ␣2 , the ␣2 phase may precipitate from the Widmanstätten ␣ phase through a hypoeutectic reaction, or from the ␤ matrix via homogeneous nucleation, or at the Widmanstätten laths/␤ matrix interfaces via heterogeneous nucleation in granular (block) morphology. However, it should be pointed out that the ␤-Ti phase is only thermodynamically stable above approximately 1100 ◦ C in the Ti–16Al system and it will transform into ␣-Ti below this temperature. However, transition metals and noble metals, i.e. V, Nb, Ta, Mo, Re, Cr and W etc., when added in sufficient concentrations, can stabilize the ␤ phase to lower temperatures, even to room temperature as demonstrated by the short-dotted arrow in Fig. 15. On the other hand, it is worthy noting that the similar misalignment of lamellar orientation has been also observed in the other Ti–Al systems with higher Al content ranging from 43 to 55 at.% such as DS Ti–Al–Mo(B) [1], Ti–Al–X (X = Mo, Re, W) [2], Ti–Al–Si [12], and binary Ti–Al [48,49]. It is known that the directional solidification process is mainly controlled by the thermal gradient and the withdrawal rate. Generally, high thermal gradient and low withdrawal rate are beneficial to the formation of DS microstructures. Nevertheless, it is obligatory to keep a mutual match between these two processing parameters. For a given axial heat transfer condition, withdrawal velocity is an important factor affecting the microstructure and the shape of the solidification front. In case of

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Fig. 13. Creep fracture surface of Ti–16Al–12Nb alloy after tested under different conditions. (a) 600 ◦ C/250 MPa/39.6 h/1.6%, (b) 600 ◦ C/460 MPa/0.3 h/2.0%, (c) 700 ◦ C/160 MPa/44 h/18.6%, and (d) 700 ◦ C/320 MPa/0.35 h/7.1%.

Fig. 14. Creep fractographs of Ti–16Al–12Nb–3Cr–1Mo alloy after tested under different conditions. (a) 600 ◦ C/210 MPa/82 h/1.3%, (b) 600 ◦ C/410 MPa/0.88 h/1.5%, (c) 700 ◦ C/180 MPa/29 h/20.9%, and (d) 700 ◦ C/350 MPa/0.39 h/6.4%.

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Table 4 True creep activation energies, Qtrue for the two alloys based on Eq. (2) after considering the temperature dependence of shear modulus. Alloy

True creep activation energy, Qtrue (kJ/mol)

Ti–16Al–12Nb Ti–16Al–12Nb–3Cr–1Mo

Low stress region I

High stress region II

316 318

312 303

4.2. Creep mechanisms

Fig. 15. Phase diagram of binary Ti–Al alloy indicating the effect of alloying element on the phase transformation and solidification process.

high withdrawal speed, the front will take the shape of concave and the favorable condition for DS will be lost [50] resulting in the misaligned microstructures along the growth direction. In the present study, comparatively lower temperature gradient and higher withdrawal velocity (60 mm/h) were employed for the DS process due to the limitations of the apparatus. In contrast, a thermal gradient as high as 400 ◦ C/cm was achieved and well-aligned lamellar microstructure was obtained in Luo et al.’s recent investigation on DS Ti–43Al–3Si alloy [12]. It is well-known that the brittleness of TiAl-based alloys at ambient temperature is closely associated with the intrinsic lattice structures of Ti3 Al (ordered hexagonal DO19 -structure) and TiAl (ordered face-centered tetragonal L10 structure), where the motion of dislocations is severely limited. Accordingly, the efforts for enhancement in the ductility of Ti–Al-based alloys via microalloying should be addressed on improving the inhomogeneous distribution of electron cloud around Ti atoms in the lattice to increase the mobility of dislocations. It is documented that the microalloying elements should occupy the substitutional positions on Al sublattice and have greater electronegativities and smaller atomic radii than Ti and/or Al, respectively [51]. As listed in Table 3 [52,53], Ti and Al atoms have an identical electronegativity value of 1.5 with a corresponding atomic radius of 0.239 and 0.236 nm. It is obvious that both Cr and Mo satisfy the aforementioned demands, and therefore, can contribute to the improvement in the inhomogeneous distribution of electron cloud around Ti atoms. As a result, the Cr- and Mo-containing Ti–Ti3 Al–Nb alloys in the present study exhibited more significantly enhanced ductility at room and high temperatures than Ti–Ti3 Al–Nb alloys. In particular, significant plastic deformation was observed at high temperatures as indicated by the fracture surfaces. Furthermore, the additions of Cr and Mo improved the mechanical properties including high temperature creep resistance via solid solution- and precipitationstrengthening effects.

Table 3 Atomic radii and electronegativities of relevant elements [52,53]. Element

Atomic radius (nm)

Electronegativity

Ti Al Nb Cr Mo

0.239 0.236 0.243 0.225 0.239

1.5 1.5 1.6 1.66 2.16

Table 4 gives the calculated true activation energy for creep in the two alloys from Eq. (2) after the temperature dependence of shear modulus was taken into account. It can be then found that the modified values decreased from 340 kJ/mol to a narrow range of 303–324 kJ/mol and fell into or were quite close to that for the self-diffusion of Ti in ␣2 -Ti3 Al (288–312 kJ/mol). Nevertheless, the significant difference in the creep exponents suggests that the different mechanisms should be operative in controlling the creep deformation at low and high stress levels. Extensive investigations have demonstrated that the creep exponents of 3–7 and Q (creep activation energy) = Qsd (self-diffusion activation energy) are indicative of a dislocation glide/climb-controlled creep mechanism [54,55]. Consequently, the similar controlling process might be operative in creep deformation for the present two alloys at 600 and 650 ◦ C in LSR and at 700 ◦ C in the full stress range. However, the higher exponents imply that some mechanisms other than conventional dislocation-controlled creep, may be operational in HSR at 600 and 650 ◦ C. The investigation by Srikant et al. [56] on creep deformation in pure-Ti demonstrated a transition in stress exponent from 7 in LSR to a higher value characterized by a power-law breakdown (PLB) in HSR. The similar result has been also obtained in Tobolová and ˇ Cadek’s study on creep of high-purity aluminum [54], where the slope of log ε˙ vs. log  curve increased and deviated from the linear relationship with stress increasing. Generally, PLB occurs at low temperature and high stress and the creep rate in this region can be usually described by [56] ε˙ = A exp

 −Q  RT

exp

 B  GT

.

(3)

However, there have still been intense disputes at the moment with regard to the scientific and reasonable mechanism operating in the high stress power-law breakdown region [57,58]. In the present study, although the alloys exhibited a high stress exponent in HSR, no deviation from the linear relationship between log ε˙ and log  was observed. Therefore, the possibility for the transition from power-law creep to PLB region can be ruled out. As stated earlier, higher creep exponents between 7 and 19 in HSR have been reported in several investigations on creep behavior of ␣2 -Ti3 Al and ␥-TiAl-based alloys [16,22,26]. In Morris and Leboeuf’s study, a stress exponent as high as 19 was obtained for creep in both lamellar Ti–48Al and duplex microstructured Ti–48Al–2Mn–2Nb alloys [22]. Such a high stress exponent has been interpreted in terms of a threshold stress existing at ␣2 /␥ or at the deformation-induced twinning interfaces that need to be overcome to initiate dislocation mobility. However, no deformation-induced twinning was observed in the present crept specimens, and moreover, deformed twinning has not been shown to always yield higher creep exponents [26]. Instead, dynamic recrystallizaion occurred as observed in the present crept specimens. In this case, the grain size decreased during creep deformation. Based on Eq. (3) where p is generally equal to 2 except for the case of grain-boundary diffusion-controlled creep for which p takes a value of 3, the experimentally measured apparent stress exponent, napp will increase to n + n, where n is the stress expo-

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nent prior to the occurrence of recrystallization and n is the increment due to dynamic recrystallization determined by [26] n = napp − n =

2 ln (d1 /d2 ) , ln (2 /1 )

(4)

where d1 and d2 denotes the average grain size before and after dynamic recrystallization taking place during the creep deformation. It can be inferred from this equation that the decrease in grain size due to dynamic recrystallization leads to an increase in the stress exponent. Nevertheless, it is difficult to quantitatively determine the dynamically recrystallized grain size since dynamic recrystallization occurred locally instead of integrally as indicated in Fig. 10b. On the other hand, the above semi-quantitative rationale considered only an average grain size in a lamellar structured system without involving the degree of recrystallization. In general, the DS microstructure is far from thermodynamical equilibrium, and therefore is unstable. This is attributed mainly to the relatively fast cooling DS process, resulting in a high energy and a high density of dislocations at grain boundaries. At high temperatures, the atomic diffusion rate increases, which helps to the occurrence and progress of dynamic recrystallization process. However, in the case of higher temperatures and stresses, the exposure time is too short to induce recrystallized grains. The similar microstructural change or instability due to dynamic recrystallization has been also observed during creep in some other lamellar TiAl-based alloys [25–27]. This implies that some lamellar structures may be inherently unstable under creep conditions. Accordingly, it is indispensable to make special efforts such as via alloying and/or heat treatment to stabilize their microstructures. 5. Conclusions (1) The alloys displayed a uniform basket-weave type and a composite-like microstructure of Ti–Ti3 Al–Nb whereas the lath orientation was not parallel to the withdrawal direction due to the limited thermal gradient and comparatively high withdrawal velocity. Nevertheless, the Mo and Cr combinative additions promoted the oriented growth, reduced the inter-lamellar spacing, and therefore remarkably enhanced the strength and ductility as well as creep resistance of the ternary Ti–16Al–12Nb alloy. (2) Transitions in creep mechanisms were noted as the stress exponent varied from a lower value of 3.5–5.3 (LSR-region I) to a higher value of 8.3–12.6 (HSR-region II) for the two alloys with increasing stress at 600 and 650 ◦ C. No transition was observed at 700 ◦ C and a single stress exponent value of 5.3–6.0 was obtained throughout the full range of applied stress from 100 to 350 MPa. The true activation energy for creep was calculated to be 303–324 kJ/mol without a transition, which fell into or was quite close to that for the self-diffusion of Ti in ␣2 -Ti3 Al (288–312 kJ/mol). The creep deformation in region I was dislocation-controlled whereas the abnormally high stress exponent in region II was associated with the effects of recrystallization owing to which the overall grain size of the initial structure was reduced. (3) The feature of fracture surfaces for the alloys was transgranular cleavage fracture at 600 ◦ C under low stress. Conversely, the specimens tested at 700 ◦ C under low stress exhibited a ductile-looking creep fracture and creep cavities were formed indicating remarkable plastic deformation during long-term creep deformation. Acknowledgements This work was supported by the Natural Science Foundation of China (No. 10672153) and “New Century Excellent Talents in Uni-

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