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Microstructure and creep behaviors of a high Nb-TiAl intermetallic compound based alloy
Materials Science & Engineering A 614 (2014) 338–346
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and creep behaviors of a high Nb-TiAl intermetallic compound based alloy Sugui Tian a,n, Qi Wang a, Huichen Yu b, Haofang Sun a, Qiuyang Li a a
School of Materials Science and Engineering, Shenyang University of Technology, Shenyang 110870, China Beijing Key Laboratory of Aeronautical Materials Testing and Evaluation, Science and Technology on Advanced High Temperature Structural Materials Laboratory, AVIC Beijing Institute of Aeronautical Materials, Beijing 100095, China b
art ic l e i nf o
a b s t r a c t
Article history: Received 24 March 2014 Received in revised form 18 May 2014 Accepted 30 June 2014 Available online 8 July 2014
By means of heat treatment, creep properties measurement and microstructure observation, an investigation has been made into the influence of heat treatment on microstructure and creep properties of the high Nb-TiAl alloy. Results show that the microstructure of as-cast high Nb-TiAl alloy consists of lamellar γ/α2 phases with various orientations. The irregular serrated boundaries with single γ phase are located in between the lamellar γ⧸α2 phases with various orientations. After solution and aging treatment, the microstructure of alloy consists of uniform and regular lamellar γ⧸α2 phases, and the regular boundaries are located in between the lamellar γ⧸α2 phases. Under the applied stress of 200 MPa at 800 1C, the creep lifetime of the as-cast high Nb-TiAl alloy is measured to be 147 h, the one of the alloy after heat treatment is measured to be 297 h. In the ranges of the applied temperatures and stresses, the creep activation energy of alloy after heart treatment is measured to be Q ¼432 kJ/mol. The deformation mechanism of alloy during creep is dislocations slipping in the lamellar γ⧸α2 phases, the creep dislocations may be decomposed to form the configuration of the partials plus stacking faults. The deformation of as-cast alloy during creep occurs mainly in the irregular serrated boundary regions with single γ phase. For the heat treated alloy, the primary/secondary slipping systems of dislocations are alternately activated during creep, which results in the bigger plastic deformation of alloy to contort the lamellar α2/γ phases. In the latter stage of creep, the cracks are firstly initiated along the boundaries parallel to the lamellar γ⧸α2 phases, and propagated along the boundaries vertical or at about 451 angles relative to the stress axis up to the occurrence of creep fracture. & 2014 Published by Elsevier B.V.
Keywords: High Nb-TiAl alloy Heat treatment Microstructure Creep Initiation and propagation of cracks Fracture mechanism
1. Introduction TiAl alloys have been widely investigated due to their excellent integrated mechanical properties [1–3], such as excellent strength, creep resistance and oxidation resistance at high temperatures [4–6]. Although the TiAl alloys display the poor ductility at room temperature [7–9], adding the elements Nb, W, Cr may improve the ductility of them at room and high temperatures [10,11]. Especially, the TiAl alloys possess a high specific strength, specific elastic module. Moreover, the density of the TiAl alloys is only half of Ni-based superalloys, and some properties of them are similar to those of Ni-based superalloys [12–14]. Therefore, the TiAl alloys are regarded as the better high-temperature structural materials with potential applications prospect for applying in aeronautics and astronautics field [15,16]. They are expected to be used for n Corresponding author. Tel.: þ 86 24 25494089 (office), þ 86 13889121677 (mobile); fax: þ86 24 25496768. E-mail address: [email protected] (S. Tian).
http://dx.doi.org/10.1016/j.msea.2014.06.117 0921-5093/& 2014 Published by Elsevier B.V.
making the hot parts in aero-engines to replace the high-density metal materials. When the components made of high Nb-TiAl alloy work for longer time at high temperatures, the creep damage is a common failure model, so that the better creep resistance of alloy at high temperatures is considered to be one of the important using criterions for preparing structure parts. At the same time, the better ductility is needed for avoiding the abrupt failure of the parts in service, which is thought to be the main evaluation criterion of high Nb-TiAl alloy for replacing the nickel-based superalloys as a weight-loss material. Microstructure of the high Nb-TiAl alloy consists of lamellar phases structure, and the mechanical and creep properties of high Nb-TiAl alloy have close relationships with their microstructure and deformation mechanisms, such as dislocations slipping, twinning and so on [17–19]. And the various microstructures of high Nb-TiAl alloy may be obtained, by different heat-treatment regimes [20,21], for displaying the different creep resistance due to the difference of the deformation mechanisms [22]. Although the
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effects of casting technologies and alloying on microstructure and mechanical properties of high Nb-TiAl alloy had been reported [23–25], the influence of heat treatment on microstructure and creep properties of high Nb-TiAl alloy is still unclear. In the paper, by means of heat treatment and creep-property measurement, combined with the microstructure observation, the influence of heat treatment on microstructure and creep properties of high Nb-TiAl alloy is investigated to provide the theory basis for promoting the development and application of high Nb-TiAl alloy.
2. Experimental procedure The high Nb-TiAl alloy was fabricated by a vacuum induction skull melting technique, and then re-melted for three times in an electric slag furnace for making the ingot of 200 mm in diameter. The nominal composition of the alloy is Ti–44Al–8Nb–0.2W–0.2B– 0.1Y alloy. The ingot of the alloy is cut into some billets with the sizes of 14 mm 40 mm 40 mm, and then some of the billets are heat treated, the heat treated regime of the billets is given as follows: 1320 1C 20 min for oil cooling, and 1250 1C 8 h for furnace cooling. After some billets of as-cast high Nb-TiAl alloy are heat treated, the as-cast and heat treated billets were cut into the specimens with cross-section of 4.5 mm 2.5 mm and gauge length of 20 mm. And uni-axial tensile creep tests were performed under constant load, in a GWT504-model creep testing machine, for measuring creep curves of the alloy at different conditions. Furthermore, in the ranges of the applied stresses and temperatures, the apparent creep active energy and apparent stress exponent of the alloy are calculated according to the data of creep curves. And the microstructures of as-cast, heat treated and creep ruptured alloys are observed under scanning electron microscopy (SEM) and transmission electron microscope (TEM), for investigating the effect of heat treatment on the microstructure and creep behavior of the alloy.
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grain, the boundaries consist of irregular serrated γ phase [26], as marked by the arrow in Fig. 1(a), and located in between the lamellar structure with different orientations. After heat treatment, the microstructure of alloy still consists of lamellar structure, as shown in Fig. 1(b), the lamellar phases within one grain are arranged along the same direction, and displaying the wavy-like configuration. The smooth grain boundaries are located in between the lamellar structures with various orientations, compared to the irregular serrated boundaries in Fig. 1(a), the boundaries in Fig. 1(b) display a smooth feature. Moreover, there are some needle-like rod phase precipitated along the boundary, the rod-like precipitate is identified as the TiB phase [27], as marked by the arrow in Fig. 1(b). XRD spectrums of TiAl–Nb alloy at different states at room temperature are measured, as shown in Fig. 2, indicating that the microstructures of the alloy at different states are mainly composed of γ-TiAl and α2-Ti3Al phases, and no rod-like TiB phase is detected in the alloy due to its small amount. But the diffraction peaks of the alloy at different states display the various intensity at special angles, this is attributed to the change of the diffraction intensity for the crystal planes of the phases. It is indicated from the XRD spectrums that the weaker intensity of the diffraction peaks appears in about 391 and 411 of 2θ angle, respectively, which corresponds to the diffraction peaks of the (111)γ and (201)a2 planes of γ and α2 phases in as-cast alloy. And the intensity of those peaks increases obvious after solution and solution þageing treatment of the alloy. Although the higher intensity of the diffraction peaks appears in the (311)γ and (222)a2 planes of γ
α − Ti3Al γ − TiAl (111) γ (201) α
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3.1. Influence of heat treatment on microstructure The microstructures of the as-cast and heat treated TiAl–Nb alloys are shown in Fig. 1, it is indicated that the microstructure of as-cast alloy consists of lamellar structure which includes the lamellar black and white phases, as shown in Fig. 1(a). The lamellar structure in Fig. 1(a) includes several lamellar congeries with parallel feature, as marked by the letters A, B and C. The lamellar structure with the same orientation in the alloy is defined as one
solution aging
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Fig. 2. XRD patterns and phases analysis of as-cast TiAl–Nb alloy after heat treated by different regimes.
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Fig. 1. Microstructures of TiAl–Nb based alloy at different states. (a) As-cast alloy, (b) heat treated alloy.
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and α2 phases in the as-cast alloy, which corresponds to the 791 of 2θ angle in the XRD spectrum, the intensity of those peaks decreases obvious after solution and solution þageing treatment of the alloy. Although the diffraction peak with middle intensity appears in the (421)a2 plane of the α2 phase for as-cast alloy, as shown in the position about 1121 of 2θ angle, the intensity of the peak increases and decreases obviously after solution and solutionþ ageing treatment of the alloy, respectively. This suggests that the volume fractions of α2-Ti3Al and γ-TiAl phases in the alloy change with heat treatment by different regimes, respectively, for example, the irregular serrated γ phase in the boundary regions of the as-cast alloy (as marked by arrow in Fig., 1a) is transformed into the lamellar γ⧸α2 phases during heat treatment, as shown in Fig. 1(b). After heat treatment, TEM images of the lamellar α2/γ phases in the TiAl–Nb alloy is shown in Fig. 3(a), which indicates that the lamellar α2/γ phases in the alloy are parallel to each other, as marked by the letters D and E, respectively. Moreover, according to SAD pattern in Fig. 3(b), the lamellar D phase is identified as γ-TiAl with Ll0 tetragonal structure, and the lamellar E phase is identified as α2-Ti3Al with D019 structure. The size of the lamellar γ⧸α2 phases in thickness is about 50–100 nm, and the lamellar γ⧸α2 phases with parallel feature are alternately arranged along the same orientation. The crystallography relationship between the lamellar γ-TiAl and α2-Ti3Al phases is identified as: (110)a2//(200)γ ̄ γ of α2-Ti3Al and γ-TiAl phases due to the and [1 1̄0]a2//[0 11] directions of incident beam for the α2-Ti3Al and γ-TiAl phases ̄ γ, respectively. Moreover, the being B ¼ [1 1̄0]a2 and B ¼ [0 11] (002)a2 plane of α2-Ti3Al phase keeps coherent interface with the (111)γ of γ-TiAl phase.
lasting time of alloy during steady state creep is about 200 h, and the creep lifetime of alloy is measured to be 297 h. Therefore, compared to the as-cast alloy, the heat treated alloy displays a better ductility in the initial period of creep and lower strain rate during steady state creep, and the bigger strain value of about 15% occurs in the alloy after crept for 297 h up to fracture. Therefore, it may be concluded that the ductility and creep resistance of the alloy may be obviously improved by means of heat treatment. Under the different conditions, the creep curves of the alloy after heat treatment are shown in Fig. 5. The creep curves of the alloy under the applied different stresses at 800 1C are shown in Fig. 5(a), indicating that, under the applied stress of 200 MPa, the alloy displays a better creep resistance and a longer creep life. When the applied stress increases to 220 MPa, the strain rate of alloy during steady-state creep is measured to be 0.023%/h, the lasting time of the alloy during steady-state creep shortens to about 120 h, and the creep lifetime of the alloy is measured to be 235 h. When the applied stress increases to 240 MPa, the strain rate of alloy during steady-state creep is measured to be 0.041%/h, the lasting time of the alloy during steady-state creep further shortens to about 50 h, and the creep lifetime of alloy is measured to be 145 h. The creep curves of the heat treated alloy under the applied stress of 200 MPa at different temperature are shown in Fig. 5(b), when the creep temperature increases to 820 1C, the strain rate of alloy during steady-state creep is measured to be 0.022%/h, the creep lifetime of the alloy is measured to be 177 h. When the creep temperature increases to 840 1C, the strain rate of the alloy during steady-state creep is measured to be 0.081%/h, the creep lifetime of the alloy is measured to be 52.5 h. This indicates that the creep feature of the alloy after heat treatment displays an obvious sensitivity to the applied temperature when the creep
3.2. Influence of heat treatment on creep properties
20 1 -- As-cast alloy 2 -- Heat treated alloy T -- 800 C σ -- 200 MPa
15
Strain, ε (%)
The creep curves of the as-cast and heat treated TiAl–Nb alloy at 800 1C/200 MPa are measured, as shown in Fig. 4. The creep curve of the as-cast alloy is marked by the number 1, which indicates that the initial strain of the as-cast alloy occurs at the moment of the applied load, and the strain rate of the alloy decreases as the creep goes on. After crept for 20 h, the creep of the alloy enters steady-state stage, the strain rate of the alloy during steady-state creep is measured to be 0.011%/h, and creep lifetime of the alloy is measured to be 147 h. After heat treatment, the creep curve of the alloy at 800 1C/ 200 MPa is marked by the number 2, indicating that, compared to the as-cast alloy, the heat treated alloy displays a bigger strain rate and strain value in the initial period of creep. After crept for 30 h, the creep of the alloy enters steady-state stage, the strain rate of alloy during steady-state creep is measured to be 0.0087%/h, the
10 2 1
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Time, (h) Fig. 4. Influence of heat treatment on creep properties of TiAl–Nb alloy.
(110)α2 (111)γ (002)α2
(112)α2
(200)γ
(311)γ
0.2μm ̄ γ. Fig. 3. After heat treatment, microstructure and SAD pattern of TiAl–Nb based alloy. (a) TEM image, (b) SAD pattern, B ¼ [1 1̄0]a2 and B ¼ [0 11]
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Fig. 5. After heat treatment, creep curves of alloy at different condition. (a) Under the applied different stresses at 800 1C, (b) under the applied stress of 200 MPa at different temperature.
-6.3 -8.8 Q = 432 kJ/mol
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Fig. 6. Dependence of the strain rate of alloy during steady-state creep on applied temperatures and stresses. (a) Strain rates and temperatures, (b) strain rates and stresses.
temperature is more than 820 1C, because the decreasing extent of the creep life for the alloy is about 230% as the creep temperature enhances from 820 1C to 840 1C. 3.3. Constitutive equation and relative parameters The transient strain of alloy occurs at the moment of the applying load at high temperatures, significant amount of dislocations are activated for slipping in the matrix of the alloy as the creep goes on, so that the density of dislocations increases to decrease the strain rate due to the effect of the strain hardening. The strain rate of the alloy maintains constant once the creep of alloy enters steady-state stage, and the strain rate of the alloy during steady-state creep may be expressed by Dorn law given as follows: Q ð1Þ ε_ ss ¼ Aσ nA exp RT where ε_ ss is the strain rate during steady-state creep, A is the constant related to material structure, σ A is the applied stress, n is the apparent stress exponent, R is the gas constant, T is the absolute temperature, and Q is the apparent creep activation energy. According to the data in the creep curves of Fig. 5, in the ranges of the applied temperatures and stresses, the strain rate of the alloy during steady-state creep are measured, and then the dependence of the strain rates of alloy during steady-state creep on the applied temperatures and stresses are expressed as
ðln ε_ ss 1=TÞ and ðln ε_ ss ln σ a Þ, as shown in Fig. 6(a) and (b), respectively. Therefore, the apparent creep activation energy and stress exponent of the alloy during steady-state creep are measured to be Q¼432 kJ/mol and n ¼4.08, respectively. According to the stress exponent, it may be deduced that the climbing of dislocations dominates the strain rate of the alloy during steadystate creep. 3.4. Deformation features of alloy during creep Under the applied stress of 200 MPa at 800 1C, the microstructure of as-cast alloy crept for 147 h up to rupture is shown in Fig. 7, which indicates that the original lamellar α2/γ phases structure still remains in the alloy, as shown in Fig. 7(a). The lamellar α2/γ phases with the same orientation are defined as one grain, the lamellar α2/γ phases with the various orientations constitute the different grains of alloy, and the boundaries are located in between the lamellar α2/γ phases with various orientations. If the region A in Fig. 7(a) is regarded as being one grain, the grain B of the lamellar α2/γ phases with another orientation passes through the grain A, which makes the grain A being separated into several fine grains. Moreover, the magnified morphology of the grain boundaries in as-cast alloy is shown in Fig. 7(b), where the boundaries with irregular serrated configuration are composed of γ phase (as marked in Fig. 7b), and located in between the lamellar α2/γ phases with various orientations. Compared to the boundary of ascast alloy in Fig. 1(a), the size of the irregular serrated boundaries in the alloy after crept up to fracture increases, as shown in Fig. 7
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γ
A
γ
B
40μm
10μm
Fig. 7. Microstructure of as-cast alloy after crept for 145 h up to fracture at 200 MPa/800 1C, (a) macro-morphology, (b) magnified morphology of boundaries.
C
D
σ
σ 20μm
20μm
30μm
Fig. 8. After heat treatment, microstructures of the alloy crept for 235 h up to rupture under the applied stress of 220 MPa at 800 1C. (a) Needle-like phase precipitated along the boundaries, (b) various space distances between the lamellar α2/γ phases appearing in different grains, (c) contorting configuration of the lamellar α2/γ phases.
(b), which is attributed to the bigger plastic deformation occurring in the region of alloy during creep. It is thought by analysis that, compare to the lamellar α2/γ phases, the single γ phase in the alloy possesses lower creep resistance and better ductility [25], therefore, the bigger plastic deformation in the alloy during creep occur in the single γ phase regions. After heat treatment, the microstructure of TiAl–Nb alloy crept for 235 h up to rupture under the applied stress of 220 MPa at 800 1C are shown in Fig. 8, the direction of the applied stress is marked by the arrow. Compared to as-cast alloy, microstructure of the heat treated alloy still consists of the regular lamellar α2/γ phases with various orientations, but the lamellar α2/γ phases display much more uniform and regular configuration, as shown in Fig. 8(a). The smooth and straight boundaries are located in between the regular and uniform lamellar α2/γ phases with various orientations, no irregular serrated single γ phase is detected in the boundary regions of alloy. The lamellar α2/γ phases with the various orientations are distributed within the various grains, some needle-like rods are precipitated along the boundary of alloy, which is identified as the TiB phase [27], as marked by the arrow in Fig. 8(a). But no needle-like precipitates are detected in another region, as shown in Fig. 8(b). Moreover, the various space distances between the lamellar α2/γ phases display within the different grains, as shown in the regions C and D, which is related
to the inclined angle of the lamellar α2/γ phases arranging relative to the normal direction of the specimen surface. The wider space distance between the lamellar α2/γ phases appears in the region C when the arranging orientation of the lamellar ones is parallel to the normal direction of the specimen surface, and the smaller space distance between the lamellar α2/γ phases is attributed to the bigger inclined angle of the lamellar ones arranging relative to the normal direction of the specimen surface, as marked in the region D of Fig. 8(b). But the lamellar α2/γ phases in the region near the fracture displays the contorting morphology, as marked by the arrow in Fig. 8(c). It is thought by analysis that the contorting configuration of the lamellar α2/γ phases is attributed to the bigger plasticity deformation occurring in the region. Especially, the primary/ secondary slipping systems of dislocations are alternately activated in the latter stage of creep, which may result in the bigger plastic deformation of alloy to contort the lamellar α2/γ phases. After heat treatment, the deformation features of the alloy crept for 297 h up to fracture at 800 1C/200 MPa is shown in Fig. 9. Some stacking faults appear in the specimen during creep, as marked by the horizontal arrow in Fig. 9(a), and some dislocations slipping in the matrix shear through the stacking fault, as marked by the inclined arrow. In another local region, the morphology of dislocations slipping in the matrix is shown in Fig. 9(b). The
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G
F 0.2μm
E
0.2μm
0.5μm
Fig. 9. After heat treatment, TEM morphologies of TiAl–Nb alloys crept for 297 h up to fracture at 800 1C/200 MPa. (a) Stacking faults, (b) dislocation slipping in matrix, (c) interaction of the stacking faults and dislocations.
dislocations with wavy-like feature slip from the left side to right side of the specimen, as marked by the shorter arrow, the inclined stacking fault stripes shear through the dislocations with wavylike feature, as marked by the longer arrow, and dislocations tangles are piled up in the region E. In another local region near the fracture, the morphology of the interaction of dislocations and stacking fault stripes is shown in Fig. 9(c). Here, some twisted dislocations are marked by the arrow, and some dislocations tangles are piled up in the region F. The stacking fault stripes are located within the region G with dislocations tangled, which is attributed to the bigger plastics deformation occurring in the region, and it is thought that the stacking fault may increase the resistance of dislocations slipping to hinder the dislocations motion. Therefore, it may be concluded that the deformation mechanism of the alloy during creep is the dislocations slipping in the matrix, the deformation dislocations slipping in matrix may be decomposed to form the configuration of the partials plus the stacking fault. And it is thought that the interaction of the deformed dislocations and stacking faults may restrain the slipping of dislocations to improve the creep resistance of alloy. 3.5. Initiation and propagation of crack during creep As creep goes on, significant amount of dislocations slipping in the matrix of the heat treated alloy are piled up in the region near the boundaries as shown in Fig. 9(b), which may cause the stress concentration due to the boundaries hindering the dislocations movement. Therefore, the cracks are firstly initiated along boundaries when the value of the stress concentration excesses the yield strength of boundary. After the alloy is crept for different time at 800 1C/240 MPa, the morphologies of the cracks initiated and propagated along boundaries are shown in Fig. 10. After the alloy is crept for 100 h, the morphology of the crack initiated along boundary is shown in Fig. 10(a), the direction of the applied stress is marked by the arrow. It can be understood from Fig. 10(a) that the crack is firstly initiated and propagated along the boundary between the lamellar α2/γ phases with different orientations, the boundary is perpendicular to the stress axis, as marked by the arrow in Fig. 10(a), which suggests that the boundaries between the lamellar α2/γ phases with different orientations is the weaker region of creep strength in alloy during creep.
Although the initiation of the crack may release the strain energy to decrease the value of the stress concentration in the region, the value of the stress concentration in the region increases again as creep goes on, which results in the propagation of the cracks in the region, so that some cavities appear in the triple junction region of the boundaries after alloy is crept up to fracture, as shown in Fig. 10(b). In another local region near the fracture, the propagation of the crack is stopped when the crack propagates to the triple junction region of the boundaries, as shown in Fig. 10(c), which may cause again the stress concentration as the creep goes on, so that some new cracks are initiated along the adjacent boundaries, as marked by the fine arrows in Fig. 10(c). After the heat treated alloy is crept up to fracture at 800 1C/ 240 MPa, in another local region, the morphologies of the cracks initiating and propagating along boundaries are shown in Fig. 11, in which the direction of the applied stress is marked by the arrow. The crack in the alloy is initiated along the boundary about 451 angles relative to the stress axis, and the propagating direction of the crack is parallel to the orientation of the lamellar α2/γ phases, as marked by the arrow in Fig. 11(a). In another local area, although the propagating direction of the crack is vertical to the orientation of the lamellar α2/γ phases, as marked by the arrow in Fig. 11(b), the propagating direction of the crack is parallel to the orientation of the lamellar α2/γ phases in another side of the crack. Therefore, it is concluded that the crack is easily initiated along the boundaries at about 451 angles relative to the stress axis since the direction bears the bigger shear stress under loading, and the initiation and propagation of crack occurs easily in the boundary parallel to the orientation of the lamellar α2/γ phases due to its weaker strength. As the creep time prolongs, the propagation of the crack along the boundary of about 451 angles relative to the stress axis is stopped when the crack propagates to the triple junction region of the boundaries, as shown in Fig. 11(c). This indicates that the boundaries with various orientations have a role of hindering the crack propagation. The observations on the microstructure indicate that the initiation and propagation of the cracks in alloy occur easily along the boundaries with straight and long features, and no initiation and propagation of cracks is detected in the regions with shorter boundary. Therefore, it may be deduced that refining grains and diminishing the size of the lamellar α2/γ phases may delay the initiation and propagation of cracks along boundary to improve the creep resistance of alloy.
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20μm
30μm
75μm
Fig. 10. Morphology of the crack initiated and propagated along boundaries of alloy after crept for different time at 800 1C/240 MPa. (a) Crack initiated along boundary of alloy crept for 100 h, (b) propagation of crack along boundary of the alloy crept up to fracture, (c) initiation and propagation of crack along various boundaries in the region near fracture.
σ
σ
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Fig. 11. Morphology of the crack initiating and propagating along the boundaries of alloy after crept up to fracture at 800 1C/240 MPa. (a) Initiation of crack along the boundary at about 451 angles relative to the stress axis, (b) propagation of crack along the boundary, (c) propagation of crack being stopped by the boundaries with various orientations.
4. Discussion 4.1. Influence of trace elements on microstructure and properties For TiAl alloys, the reaction of the trace B with the element Ti may occurs to form the fine TiB phase, the one acts as the particles of the neterogeny nucleation, which may both promote the uniform nucleation of α2 phase and refine the grain size of TiAl alloy. Therefore, the volume fraction of α2 phase in the alloy increases with adding the trace B to decrease the volume fraction of γ phase due to the effect of the trace B [28]. Particularly, the microstructure with α⧸γ lamellar feature in TiAl alloy possesses an obvious anisotropy feature during deformation, the deformed dislocations are easily activated in the γ phase with lower strength [25] under the action of the applied loading at high temperature. Therefore, the alloy with higher volume fraction of γ phase during creep displays the higher strain rate and shorter creep life, but the creep resistance of the alloy may be obviously improved when the volume fraction of γ phase is less than 12% [29].
Adding element Y may both enhance the resistance oxidation and refine the grain sizes of TiAl alloy, as well as improve the microstructure of as-cast TiAl alloys [30]. Especially, when the elements B and Y are added to TiAl alloy together, the better effect of the refining grains may be obtained [31]. On the one hand, adding the element W may promote the precipitation of the initial β phase in the TiAl alloy during solidification, therefore, compared to the TiAl–Cr alloy, TiAl–W alloy reserves the higher volume fraction of β phase [32], which may promote the occurrence of the phases transformation given as follows [33,34]: L-Lþ β-αþβ-α-αþ γ-α2 þγ. On the other hand, the element W may delay the process of the deformation dynamics being dominated by diffusion mechanisms, due to its lower diffusion coefficient, which may enhance the creep properties and the transformation temperature of the toughness/brittleness of TiAl alloys [28]. Therefore, the yield strength of TiAl alloy at 871 1C may enhance from 298 MPa to 374 MPa due to adding 0.2% W element [29]. It is indicated [35] that the creep life of high NbTiAl alloy at 800 1C/200 MPa is measured to be 57 h when the 2 wt
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% Cr and 0.15 wt% B elements add to the alloy. While the creep life of the high Nb-TiAl alloy at 800 1C/200 MPa is measured to be 147 h when the 0.2 wt% W, 0.2 wt% B and 0.1 wt% Y elements add to the alloy. After heat treatment, the creep life of the alloy enhances to 297 h, as shown in Fig. 4, which indicates that the traces W, B and Y may obviously enhance the creep life of the high Nb-TiAl alloy. 4.2. Effect of boundary configuration on the creep resistance The grain boundaries in as-cast high Nb-TiAl alloy display the irregular serrated configuration, as marked by the arrow in Fig. 1 (a), the boundaries consist of the single γ phase which has a weaker strength and better ductility [25]. Although the smaller strain occurs in as-cast alloy during creep at high-temperature, for example, the strain value of the as-cast alloy crept up to fracture is only about 7%, as shown in as shown in Fig. 4, the bigger plastic deformation occurs in the boundary regions with single γ phase, as shown in Fig. 7(b). This indicates that, compared to the lamellar α2/γ phases, a lower strength appears in the irregular serrated boundary regions with single γ phase in the alloy. Therefore, the strain of the alloy during creep occurs mainly in the irregular serrated boundary regions with single γ phase. Namely, the irregular serrated boundaries with single γ phase are thought to be the weak region of the creep strength during creep of the ascast high Nb-TiAl alloy. After heat treatment, the irregular serrated boundaries with single γ phase in the alloy disappear, and the grain boundary displays the smooth and straight features, as shown in Fig. 1(b), which may improve the creep resistance of alloy. Therefore, the creep life of the alloy after heat treatment increases from 147 h to 297 h under the applied stress of 200 MPa at 800 1C, as shown in Fig. 4, the increasing extent of the creep life is about 102%, which indicates that eliminating irregular serrated boundaries with single γ phase may improve the creep properties of the alloy to a great extent. The microstructure of both as cast and heat treated high NbTiAl alloy consists of lamellar α2/γ phases, and the lamellar α2/γ phases within grains display the wavy-like feature, as shown in Fig. 1(a) and (b). But after crept up to fracture, the lamellar α2/γ phases with wavy-like feature in alloys have transformed into the regular and straight configuration, as shown in Figs. 7(a) and 8, which indicates that the configuration transformation of the lamellar α2/γ phases occurs during creep at high temperatures due to the diffusion of the elements. 4.3. Analysis on initiation and propagation of crack during creep Although the irregular serrated boundaries with single γ phase may be eliminated, by means of heat treatment at high temperature, to enhance the creep properties of the alloy to a great extent, the crack in the alloy during creep is firstly initiated and propagated along the boundary, which suggests that the boundaries is still the weaker region of the creep strength in the alloy during creep. As creep goes on, significant amount of dislocations are activated in the alloy for piling up in the region near the boundaries, as shown in Fig. 9(b), which may cause the stress concentration to promote the initiation and propagation of the crack along the boundary up to the occurrence of creep fracture. Actually, the initiation of the crack may releases the stress concentration in the region. But the stress concentration may occur again in the region near the crack as creep goes on, which results in the propagation of the crack along the boundary. Especially, the grain boundaries perpendicular or at 451 angles relative to the stress axis bear the bigger shearing stress during creep, which causes easily the initiation and propagation of the
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cracks along the boundaries, as shown in Figs. 10 and 11. Although the boundaries with various orientations may stop the propagation of the crack when the one propagates to the triple junction region of the boundaries, the stress concentration occurs again in the region as the creep goes on, which may promote the initiation and propagation of the cracks occurring in the adjacent boundaries, as shown in Fig. 10(b) and (c). As the propagation of the multi-cracks in different cross-sections goes on, the tearing edges and secondary cracks are formed at the tip region of the primary crack along the direction with bigger shearing stress, which results in the gradual connection of the multi-cracks in different crosssections, by the tearing edges and secondary cleavage plane, up to the occurrence of creep fracture [36]. This is thought to be the fracture mechanism of the alloy during creep.
5. Conclusions (1) The microstructure of the as-cast high Nb-TiAl alloy consists of the lamellar γ⧸α2 phases with various orientations, and the irregular serrated boundaries with single γ phase are located in between the lamellar ones. After heat treatment, the irregular serrated single γ phase disappears in the boundary regions, and the regular boundaries are formed in between the lamellar α2/γ phases with various orientations. (2) Under the applied stress of 200 MPa at 800 1C, the creep life of as-cast high Nb-TiAl alloy is measured to be 145 h, and the one of the alloy after heat treated is enhanced to 297 h due to eliminating the irregular serrated boundaries with single γ phase. The creep activation energy of the heat treated alloy is measured to be Q¼ 432 kJ/mol in the ranges of the applied temperatures and stresses. (3) The deformation mechanism of alloy during creep is dislocations slipping in the matrix, the dislocations slipping in the matrix may be decomposed to form the configuration of the partials plus the stacking faults. The interaction of the deformation dislocations and stacking faults may hinder dislocations movement to improve the creep resistance of alloy. (4) The strain of as-cast alloy during creep occurs mainly in the irregular serrated boundary regions with single γ phase. After the alloy is heat treated, the primary/secondary slipping systems of dislocations are alternately activated in the latter stage of creep, which may result in the bigger plastic deformation for contorting the lamellar α2/γ phases in alloy. (5) The cracks in the alloys during creep are firstly initiated and propagated in the boundaries, which are parallel to the lamellar γ⧸α phases, along the direction perpendicular or at about 451 angles relative to the stress axis up to the occurrence of creep fracture, which is thought to be fracture mechanism of the alloy during creep.
Acknowledgments Sponsorship of this research by the National Basic Research Program of China under Grant no. 2011CB605506 is gratefully acknowledged. References [1] X.F. Ding, J.P. Lin, L.Q. Zhang, Y.Q. Su, G.L. Chen, Acta Mater. 60 (2012) 498–506. [2] G. Hénaff, A.L. Gloanec, Intermetallics 13 (2005) 543–558.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and creep behaviors of a high Nb-TiAl intermetallic compound based alloy Sugui Tian a,n, Qi Wang a, Huichen Yu b, Haofang Sun a, Qiuyang Li a a
School of Materials Science and Engineering, Shenyang University of Technology, Shenyang 110870, China Beijing Key Laboratory of Aeronautical Materials Testing and Evaluation, Science and Technology on Advanced High Temperature Structural Materials Laboratory, AVIC Beijing Institute of Aeronautical Materials, Beijing 100095, China b
art ic l e i nf o
a b s t r a c t
Article history: Received 24 March 2014 Received in revised form 18 May 2014 Accepted 30 June 2014 Available online 8 July 2014
By means of heat treatment, creep properties measurement and microstructure observation, an investigation has been made into the influence of heat treatment on microstructure and creep properties of the high Nb-TiAl alloy. Results show that the microstructure of as-cast high Nb-TiAl alloy consists of lamellar γ/α2 phases with various orientations. The irregular serrated boundaries with single γ phase are located in between the lamellar γ⧸α2 phases with various orientations. After solution and aging treatment, the microstructure of alloy consists of uniform and regular lamellar γ⧸α2 phases, and the regular boundaries are located in between the lamellar γ⧸α2 phases. Under the applied stress of 200 MPa at 800 1C, the creep lifetime of the as-cast high Nb-TiAl alloy is measured to be 147 h, the one of the alloy after heat treatment is measured to be 297 h. In the ranges of the applied temperatures and stresses, the creep activation energy of alloy after heart treatment is measured to be Q ¼432 kJ/mol. The deformation mechanism of alloy during creep is dislocations slipping in the lamellar γ⧸α2 phases, the creep dislocations may be decomposed to form the configuration of the partials plus stacking faults. The deformation of as-cast alloy during creep occurs mainly in the irregular serrated boundary regions with single γ phase. For the heat treated alloy, the primary/secondary slipping systems of dislocations are alternately activated during creep, which results in the bigger plastic deformation of alloy to contort the lamellar α2/γ phases. In the latter stage of creep, the cracks are firstly initiated along the boundaries parallel to the lamellar γ⧸α2 phases, and propagated along the boundaries vertical or at about 451 angles relative to the stress axis up to the occurrence of creep fracture. & 2014 Published by Elsevier B.V.
Keywords: High Nb-TiAl alloy Heat treatment Microstructure Creep Initiation and propagation of cracks Fracture mechanism
1. Introduction TiAl alloys have been widely investigated due to their excellent integrated mechanical properties [1–3], such as excellent strength, creep resistance and oxidation resistance at high temperatures [4–6]. Although the TiAl alloys display the poor ductility at room temperature [7–9], adding the elements Nb, W, Cr may improve the ductility of them at room and high temperatures [10,11]. Especially, the TiAl alloys possess a high specific strength, specific elastic module. Moreover, the density of the TiAl alloys is only half of Ni-based superalloys, and some properties of them are similar to those of Ni-based superalloys [12–14]. Therefore, the TiAl alloys are regarded as the better high-temperature structural materials with potential applications prospect for applying in aeronautics and astronautics field [15,16]. They are expected to be used for n Corresponding author. Tel.: þ 86 24 25494089 (office), þ 86 13889121677 (mobile); fax: þ86 24 25496768. E-mail address: [email protected] (S. Tian).
http://dx.doi.org/10.1016/j.msea.2014.06.117 0921-5093/& 2014 Published by Elsevier B.V.
making the hot parts in aero-engines to replace the high-density metal materials. When the components made of high Nb-TiAl alloy work for longer time at high temperatures, the creep damage is a common failure model, so that the better creep resistance of alloy at high temperatures is considered to be one of the important using criterions for preparing structure parts. At the same time, the better ductility is needed for avoiding the abrupt failure of the parts in service, which is thought to be the main evaluation criterion of high Nb-TiAl alloy for replacing the nickel-based superalloys as a weight-loss material. Microstructure of the high Nb-TiAl alloy consists of lamellar phases structure, and the mechanical and creep properties of high Nb-TiAl alloy have close relationships with their microstructure and deformation mechanisms, such as dislocations slipping, twinning and so on [17–19]. And the various microstructures of high Nb-TiAl alloy may be obtained, by different heat-treatment regimes [20,21], for displaying the different creep resistance due to the difference of the deformation mechanisms [22]. Although the
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effects of casting technologies and alloying on microstructure and mechanical properties of high Nb-TiAl alloy had been reported [23–25], the influence of heat treatment on microstructure and creep properties of high Nb-TiAl alloy is still unclear. In the paper, by means of heat treatment and creep-property measurement, combined with the microstructure observation, the influence of heat treatment on microstructure and creep properties of high Nb-TiAl alloy is investigated to provide the theory basis for promoting the development and application of high Nb-TiAl alloy.
2. Experimental procedure The high Nb-TiAl alloy was fabricated by a vacuum induction skull melting technique, and then re-melted for three times in an electric slag furnace for making the ingot of 200 mm in diameter. The nominal composition of the alloy is Ti–44Al–8Nb–0.2W–0.2B– 0.1Y alloy. The ingot of the alloy is cut into some billets with the sizes of 14 mm 40 mm 40 mm, and then some of the billets are heat treated, the heat treated regime of the billets is given as follows: 1320 1C 20 min for oil cooling, and 1250 1C 8 h for furnace cooling. After some billets of as-cast high Nb-TiAl alloy are heat treated, the as-cast and heat treated billets were cut into the specimens with cross-section of 4.5 mm 2.5 mm and gauge length of 20 mm. And uni-axial tensile creep tests were performed under constant load, in a GWT504-model creep testing machine, for measuring creep curves of the alloy at different conditions. Furthermore, in the ranges of the applied stresses and temperatures, the apparent creep active energy and apparent stress exponent of the alloy are calculated according to the data of creep curves. And the microstructures of as-cast, heat treated and creep ruptured alloys are observed under scanning electron microscopy (SEM) and transmission electron microscope (TEM), for investigating the effect of heat treatment on the microstructure and creep behavior of the alloy.
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grain, the boundaries consist of irregular serrated γ phase [26], as marked by the arrow in Fig. 1(a), and located in between the lamellar structure with different orientations. After heat treatment, the microstructure of alloy still consists of lamellar structure, as shown in Fig. 1(b), the lamellar phases within one grain are arranged along the same direction, and displaying the wavy-like configuration. The smooth grain boundaries are located in between the lamellar structures with various orientations, compared to the irregular serrated boundaries in Fig. 1(a), the boundaries in Fig. 1(b) display a smooth feature. Moreover, there are some needle-like rod phase precipitated along the boundary, the rod-like precipitate is identified as the TiB phase [27], as marked by the arrow in Fig. 1(b). XRD spectrums of TiAl–Nb alloy at different states at room temperature are measured, as shown in Fig. 2, indicating that the microstructures of the alloy at different states are mainly composed of γ-TiAl and α2-Ti3Al phases, and no rod-like TiB phase is detected in the alloy due to its small amount. But the diffraction peaks of the alloy at different states display the various intensity at special angles, this is attributed to the change of the diffraction intensity for the crystal planes of the phases. It is indicated from the XRD spectrums that the weaker intensity of the diffraction peaks appears in about 391 and 411 of 2θ angle, respectively, which corresponds to the diffraction peaks of the (111)γ and (201)a2 planes of γ and α2 phases in as-cast alloy. And the intensity of those peaks increases obvious after solution and solution þageing treatment of the alloy. Although the higher intensity of the diffraction peaks appears in the (311)γ and (222)a2 planes of γ
α − Ti3Al γ − TiAl (111) γ (201) α
(311) (202)α
γ
(222)α
(421) α
As-cast
3. Experimental results and analysis γ
solution
3.1. Influence of heat treatment on microstructure The microstructures of the as-cast and heat treated TiAl–Nb alloys are shown in Fig. 1, it is indicated that the microstructure of as-cast alloy consists of lamellar structure which includes the lamellar black and white phases, as shown in Fig. 1(a). The lamellar structure in Fig. 1(a) includes several lamellar congeries with parallel feature, as marked by the letters A, B and C. The lamellar structure with the same orientation in the alloy is defined as one
solution aging
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20
40
60
80
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Fig. 2. XRD patterns and phases analysis of as-cast TiAl–Nb alloy after heat treated by different regimes.
A B C
5μm
100
2θ
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Fig. 1. Microstructures of TiAl–Nb based alloy at different states. (a) As-cast alloy, (b) heat treated alloy.
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and α2 phases in the as-cast alloy, which corresponds to the 791 of 2θ angle in the XRD spectrum, the intensity of those peaks decreases obvious after solution and solution þageing treatment of the alloy. Although the diffraction peak with middle intensity appears in the (421)a2 plane of the α2 phase for as-cast alloy, as shown in the position about 1121 of 2θ angle, the intensity of the peak increases and decreases obviously after solution and solutionþ ageing treatment of the alloy, respectively. This suggests that the volume fractions of α2-Ti3Al and γ-TiAl phases in the alloy change with heat treatment by different regimes, respectively, for example, the irregular serrated γ phase in the boundary regions of the as-cast alloy (as marked by arrow in Fig., 1a) is transformed into the lamellar γ⧸α2 phases during heat treatment, as shown in Fig. 1(b). After heat treatment, TEM images of the lamellar α2/γ phases in the TiAl–Nb alloy is shown in Fig. 3(a), which indicates that the lamellar α2/γ phases in the alloy are parallel to each other, as marked by the letters D and E, respectively. Moreover, according to SAD pattern in Fig. 3(b), the lamellar D phase is identified as γ-TiAl with Ll0 tetragonal structure, and the lamellar E phase is identified as α2-Ti3Al with D019 structure. The size of the lamellar γ⧸α2 phases in thickness is about 50–100 nm, and the lamellar γ⧸α2 phases with parallel feature are alternately arranged along the same orientation. The crystallography relationship between the lamellar γ-TiAl and α2-Ti3Al phases is identified as: (110)a2//(200)γ ̄ γ of α2-Ti3Al and γ-TiAl phases due to the and [1 1̄0]a2//[0 11] directions of incident beam for the α2-Ti3Al and γ-TiAl phases ̄ γ, respectively. Moreover, the being B ¼ [1 1̄0]a2 and B ¼ [0 11] (002)a2 plane of α2-Ti3Al phase keeps coherent interface with the (111)γ of γ-TiAl phase.
lasting time of alloy during steady state creep is about 200 h, and the creep lifetime of alloy is measured to be 297 h. Therefore, compared to the as-cast alloy, the heat treated alloy displays a better ductility in the initial period of creep and lower strain rate during steady state creep, and the bigger strain value of about 15% occurs in the alloy after crept for 297 h up to fracture. Therefore, it may be concluded that the ductility and creep resistance of the alloy may be obviously improved by means of heat treatment. Under the different conditions, the creep curves of the alloy after heat treatment are shown in Fig. 5. The creep curves of the alloy under the applied different stresses at 800 1C are shown in Fig. 5(a), indicating that, under the applied stress of 200 MPa, the alloy displays a better creep resistance and a longer creep life. When the applied stress increases to 220 MPa, the strain rate of alloy during steady-state creep is measured to be 0.023%/h, the lasting time of the alloy during steady-state creep shortens to about 120 h, and the creep lifetime of the alloy is measured to be 235 h. When the applied stress increases to 240 MPa, the strain rate of alloy during steady-state creep is measured to be 0.041%/h, the lasting time of the alloy during steady-state creep further shortens to about 50 h, and the creep lifetime of alloy is measured to be 145 h. The creep curves of the heat treated alloy under the applied stress of 200 MPa at different temperature are shown in Fig. 5(b), when the creep temperature increases to 820 1C, the strain rate of alloy during steady-state creep is measured to be 0.022%/h, the creep lifetime of the alloy is measured to be 177 h. When the creep temperature increases to 840 1C, the strain rate of the alloy during steady-state creep is measured to be 0.081%/h, the creep lifetime of the alloy is measured to be 52.5 h. This indicates that the creep feature of the alloy after heat treatment displays an obvious sensitivity to the applied temperature when the creep
3.2. Influence of heat treatment on creep properties
20 1 -- As-cast alloy 2 -- Heat treated alloy T -- 800 C σ -- 200 MPa
15
Strain, ε (%)
The creep curves of the as-cast and heat treated TiAl–Nb alloy at 800 1C/200 MPa are measured, as shown in Fig. 4. The creep curve of the as-cast alloy is marked by the number 1, which indicates that the initial strain of the as-cast alloy occurs at the moment of the applied load, and the strain rate of the alloy decreases as the creep goes on. After crept for 20 h, the creep of the alloy enters steady-state stage, the strain rate of the alloy during steady-state creep is measured to be 0.011%/h, and creep lifetime of the alloy is measured to be 147 h. After heat treatment, the creep curve of the alloy at 800 1C/ 200 MPa is marked by the number 2, indicating that, compared to the as-cast alloy, the heat treated alloy displays a bigger strain rate and strain value in the initial period of creep. After crept for 30 h, the creep of the alloy enters steady-state stage, the strain rate of alloy during steady-state creep is measured to be 0.0087%/h, the
10 2 1
5
0
0
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Time, (h) Fig. 4. Influence of heat treatment on creep properties of TiAl–Nb alloy.
(110)α2 (111)γ (002)α2
(112)α2
(200)γ
(311)γ
0.2μm ̄ γ. Fig. 3. After heat treatment, microstructure and SAD pattern of TiAl–Nb based alloy. (a) TEM image, (b) SAD pattern, B ¼ [1 1̄0]a2 and B ¼ [0 11]
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25
25
Strain, ε (%)
Strain, ε (%)
1 -- 800 C
1 -- 200MPa 2 -- 220MPa 3 -- 240MPa
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T -- 800 C
15 10
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2 -- 820 C
15
3 -- 840 C σ -- 200MPa
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Time, (h)
Time, (h)
Fig. 5. After heat treatment, creep curves of alloy at different condition. (a) Under the applied different stresses at 800 1C, (b) under the applied stress of 200 MPa at different temperature.
-6.3 -8.8 Q = 432 kJ/mol
-6.6
-9.6
-6.9
ln
ln
-10.4 -11.2 -12.0 -12.8 0.90
n = 4.08
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0.95
1.00
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1.15
1/T, 10-3K-1
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19.28
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In σ
Fig. 6. Dependence of the strain rate of alloy during steady-state creep on applied temperatures and stresses. (a) Strain rates and temperatures, (b) strain rates and stresses.
temperature is more than 820 1C, because the decreasing extent of the creep life for the alloy is about 230% as the creep temperature enhances from 820 1C to 840 1C. 3.3. Constitutive equation and relative parameters The transient strain of alloy occurs at the moment of the applying load at high temperatures, significant amount of dislocations are activated for slipping in the matrix of the alloy as the creep goes on, so that the density of dislocations increases to decrease the strain rate due to the effect of the strain hardening. The strain rate of the alloy maintains constant once the creep of alloy enters steady-state stage, and the strain rate of the alloy during steady-state creep may be expressed by Dorn law given as follows: Q ð1Þ ε_ ss ¼ Aσ nA exp RT where ε_ ss is the strain rate during steady-state creep, A is the constant related to material structure, σ A is the applied stress, n is the apparent stress exponent, R is the gas constant, T is the absolute temperature, and Q is the apparent creep activation energy. According to the data in the creep curves of Fig. 5, in the ranges of the applied temperatures and stresses, the strain rate of the alloy during steady-state creep are measured, and then the dependence of the strain rates of alloy during steady-state creep on the applied temperatures and stresses are expressed as
ðln ε_ ss 1=TÞ and ðln ε_ ss ln σ a Þ, as shown in Fig. 6(a) and (b), respectively. Therefore, the apparent creep activation energy and stress exponent of the alloy during steady-state creep are measured to be Q¼432 kJ/mol and n ¼4.08, respectively. According to the stress exponent, it may be deduced that the climbing of dislocations dominates the strain rate of the alloy during steadystate creep. 3.4. Deformation features of alloy during creep Under the applied stress of 200 MPa at 800 1C, the microstructure of as-cast alloy crept for 147 h up to rupture is shown in Fig. 7, which indicates that the original lamellar α2/γ phases structure still remains in the alloy, as shown in Fig. 7(a). The lamellar α2/γ phases with the same orientation are defined as one grain, the lamellar α2/γ phases with the various orientations constitute the different grains of alloy, and the boundaries are located in between the lamellar α2/γ phases with various orientations. If the region A in Fig. 7(a) is regarded as being one grain, the grain B of the lamellar α2/γ phases with another orientation passes through the grain A, which makes the grain A being separated into several fine grains. Moreover, the magnified morphology of the grain boundaries in as-cast alloy is shown in Fig. 7(b), where the boundaries with irregular serrated configuration are composed of γ phase (as marked in Fig. 7b), and located in between the lamellar α2/γ phases with various orientations. Compared to the boundary of ascast alloy in Fig. 1(a), the size of the irregular serrated boundaries in the alloy after crept up to fracture increases, as shown in Fig. 7
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γ
A
γ
B
40μm
10μm
Fig. 7. Microstructure of as-cast alloy after crept for 145 h up to fracture at 200 MPa/800 1C, (a) macro-morphology, (b) magnified morphology of boundaries.
C
D
σ
σ 20μm
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Fig. 8. After heat treatment, microstructures of the alloy crept for 235 h up to rupture under the applied stress of 220 MPa at 800 1C. (a) Needle-like phase precipitated along the boundaries, (b) various space distances between the lamellar α2/γ phases appearing in different grains, (c) contorting configuration of the lamellar α2/γ phases.
(b), which is attributed to the bigger plastic deformation occurring in the region of alloy during creep. It is thought by analysis that, compare to the lamellar α2/γ phases, the single γ phase in the alloy possesses lower creep resistance and better ductility [25], therefore, the bigger plastic deformation in the alloy during creep occur in the single γ phase regions. After heat treatment, the microstructure of TiAl–Nb alloy crept for 235 h up to rupture under the applied stress of 220 MPa at 800 1C are shown in Fig. 8, the direction of the applied stress is marked by the arrow. Compared to as-cast alloy, microstructure of the heat treated alloy still consists of the regular lamellar α2/γ phases with various orientations, but the lamellar α2/γ phases display much more uniform and regular configuration, as shown in Fig. 8(a). The smooth and straight boundaries are located in between the regular and uniform lamellar α2/γ phases with various orientations, no irregular serrated single γ phase is detected in the boundary regions of alloy. The lamellar α2/γ phases with the various orientations are distributed within the various grains, some needle-like rods are precipitated along the boundary of alloy, which is identified as the TiB phase [27], as marked by the arrow in Fig. 8(a). But no needle-like precipitates are detected in another region, as shown in Fig. 8(b). Moreover, the various space distances between the lamellar α2/γ phases display within the different grains, as shown in the regions C and D, which is related
to the inclined angle of the lamellar α2/γ phases arranging relative to the normal direction of the specimen surface. The wider space distance between the lamellar α2/γ phases appears in the region C when the arranging orientation of the lamellar ones is parallel to the normal direction of the specimen surface, and the smaller space distance between the lamellar α2/γ phases is attributed to the bigger inclined angle of the lamellar ones arranging relative to the normal direction of the specimen surface, as marked in the region D of Fig. 8(b). But the lamellar α2/γ phases in the region near the fracture displays the contorting morphology, as marked by the arrow in Fig. 8(c). It is thought by analysis that the contorting configuration of the lamellar α2/γ phases is attributed to the bigger plasticity deformation occurring in the region. Especially, the primary/ secondary slipping systems of dislocations are alternately activated in the latter stage of creep, which may result in the bigger plastic deformation of alloy to contort the lamellar α2/γ phases. After heat treatment, the deformation features of the alloy crept for 297 h up to fracture at 800 1C/200 MPa is shown in Fig. 9. Some stacking faults appear in the specimen during creep, as marked by the horizontal arrow in Fig. 9(a), and some dislocations slipping in the matrix shear through the stacking fault, as marked by the inclined arrow. In another local region, the morphology of dislocations slipping in the matrix is shown in Fig. 9(b). The
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G
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Fig. 9. After heat treatment, TEM morphologies of TiAl–Nb alloys crept for 297 h up to fracture at 800 1C/200 MPa. (a) Stacking faults, (b) dislocation slipping in matrix, (c) interaction of the stacking faults and dislocations.
dislocations with wavy-like feature slip from the left side to right side of the specimen, as marked by the shorter arrow, the inclined stacking fault stripes shear through the dislocations with wavylike feature, as marked by the longer arrow, and dislocations tangles are piled up in the region E. In another local region near the fracture, the morphology of the interaction of dislocations and stacking fault stripes is shown in Fig. 9(c). Here, some twisted dislocations are marked by the arrow, and some dislocations tangles are piled up in the region F. The stacking fault stripes are located within the region G with dislocations tangled, which is attributed to the bigger plastics deformation occurring in the region, and it is thought that the stacking fault may increase the resistance of dislocations slipping to hinder the dislocations motion. Therefore, it may be concluded that the deformation mechanism of the alloy during creep is the dislocations slipping in the matrix, the deformation dislocations slipping in matrix may be decomposed to form the configuration of the partials plus the stacking fault. And it is thought that the interaction of the deformed dislocations and stacking faults may restrain the slipping of dislocations to improve the creep resistance of alloy. 3.5. Initiation and propagation of crack during creep As creep goes on, significant amount of dislocations slipping in the matrix of the heat treated alloy are piled up in the region near the boundaries as shown in Fig. 9(b), which may cause the stress concentration due to the boundaries hindering the dislocations movement. Therefore, the cracks are firstly initiated along boundaries when the value of the stress concentration excesses the yield strength of boundary. After the alloy is crept for different time at 800 1C/240 MPa, the morphologies of the cracks initiated and propagated along boundaries are shown in Fig. 10. After the alloy is crept for 100 h, the morphology of the crack initiated along boundary is shown in Fig. 10(a), the direction of the applied stress is marked by the arrow. It can be understood from Fig. 10(a) that the crack is firstly initiated and propagated along the boundary between the lamellar α2/γ phases with different orientations, the boundary is perpendicular to the stress axis, as marked by the arrow in Fig. 10(a), which suggests that the boundaries between the lamellar α2/γ phases with different orientations is the weaker region of creep strength in alloy during creep.
Although the initiation of the crack may release the strain energy to decrease the value of the stress concentration in the region, the value of the stress concentration in the region increases again as creep goes on, which results in the propagation of the cracks in the region, so that some cavities appear in the triple junction region of the boundaries after alloy is crept up to fracture, as shown in Fig. 10(b). In another local region near the fracture, the propagation of the crack is stopped when the crack propagates to the triple junction region of the boundaries, as shown in Fig. 10(c), which may cause again the stress concentration as the creep goes on, so that some new cracks are initiated along the adjacent boundaries, as marked by the fine arrows in Fig. 10(c). After the heat treated alloy is crept up to fracture at 800 1C/ 240 MPa, in another local region, the morphologies of the cracks initiating and propagating along boundaries are shown in Fig. 11, in which the direction of the applied stress is marked by the arrow. The crack in the alloy is initiated along the boundary about 451 angles relative to the stress axis, and the propagating direction of the crack is parallel to the orientation of the lamellar α2/γ phases, as marked by the arrow in Fig. 11(a). In another local area, although the propagating direction of the crack is vertical to the orientation of the lamellar α2/γ phases, as marked by the arrow in Fig. 11(b), the propagating direction of the crack is parallel to the orientation of the lamellar α2/γ phases in another side of the crack. Therefore, it is concluded that the crack is easily initiated along the boundaries at about 451 angles relative to the stress axis since the direction bears the bigger shear stress under loading, and the initiation and propagation of crack occurs easily in the boundary parallel to the orientation of the lamellar α2/γ phases due to its weaker strength. As the creep time prolongs, the propagation of the crack along the boundary of about 451 angles relative to the stress axis is stopped when the crack propagates to the triple junction region of the boundaries, as shown in Fig. 11(c). This indicates that the boundaries with various orientations have a role of hindering the crack propagation. The observations on the microstructure indicate that the initiation and propagation of the cracks in alloy occur easily along the boundaries with straight and long features, and no initiation and propagation of cracks is detected in the regions with shorter boundary. Therefore, it may be deduced that refining grains and diminishing the size of the lamellar α2/γ phases may delay the initiation and propagation of cracks along boundary to improve the creep resistance of alloy.
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Fig. 10. Morphology of the crack initiated and propagated along boundaries of alloy after crept for different time at 800 1C/240 MPa. (a) Crack initiated along boundary of alloy crept for 100 h, (b) propagation of crack along boundary of the alloy crept up to fracture, (c) initiation and propagation of crack along various boundaries in the region near fracture.
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Fig. 11. Morphology of the crack initiating and propagating along the boundaries of alloy after crept up to fracture at 800 1C/240 MPa. (a) Initiation of crack along the boundary at about 451 angles relative to the stress axis, (b) propagation of crack along the boundary, (c) propagation of crack being stopped by the boundaries with various orientations.
4. Discussion 4.1. Influence of trace elements on microstructure and properties For TiAl alloys, the reaction of the trace B with the element Ti may occurs to form the fine TiB phase, the one acts as the particles of the neterogeny nucleation, which may both promote the uniform nucleation of α2 phase and refine the grain size of TiAl alloy. Therefore, the volume fraction of α2 phase in the alloy increases with adding the trace B to decrease the volume fraction of γ phase due to the effect of the trace B [28]. Particularly, the microstructure with α⧸γ lamellar feature in TiAl alloy possesses an obvious anisotropy feature during deformation, the deformed dislocations are easily activated in the γ phase with lower strength [25] under the action of the applied loading at high temperature. Therefore, the alloy with higher volume fraction of γ phase during creep displays the higher strain rate and shorter creep life, but the creep resistance of the alloy may be obviously improved when the volume fraction of γ phase is less than 12% [29].
Adding element Y may both enhance the resistance oxidation and refine the grain sizes of TiAl alloy, as well as improve the microstructure of as-cast TiAl alloys [30]. Especially, when the elements B and Y are added to TiAl alloy together, the better effect of the refining grains may be obtained [31]. On the one hand, adding the element W may promote the precipitation of the initial β phase in the TiAl alloy during solidification, therefore, compared to the TiAl–Cr alloy, TiAl–W alloy reserves the higher volume fraction of β phase [32], which may promote the occurrence of the phases transformation given as follows [33,34]: L-Lþ β-αþβ-α-αþ γ-α2 þγ. On the other hand, the element W may delay the process of the deformation dynamics being dominated by diffusion mechanisms, due to its lower diffusion coefficient, which may enhance the creep properties and the transformation temperature of the toughness/brittleness of TiAl alloys [28]. Therefore, the yield strength of TiAl alloy at 871 1C may enhance from 298 MPa to 374 MPa due to adding 0.2% W element [29]. It is indicated [35] that the creep life of high NbTiAl alloy at 800 1C/200 MPa is measured to be 57 h when the 2 wt
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% Cr and 0.15 wt% B elements add to the alloy. While the creep life of the high Nb-TiAl alloy at 800 1C/200 MPa is measured to be 147 h when the 0.2 wt% W, 0.2 wt% B and 0.1 wt% Y elements add to the alloy. After heat treatment, the creep life of the alloy enhances to 297 h, as shown in Fig. 4, which indicates that the traces W, B and Y may obviously enhance the creep life of the high Nb-TiAl alloy. 4.2. Effect of boundary configuration on the creep resistance The grain boundaries in as-cast high Nb-TiAl alloy display the irregular serrated configuration, as marked by the arrow in Fig. 1 (a), the boundaries consist of the single γ phase which has a weaker strength and better ductility [25]. Although the smaller strain occurs in as-cast alloy during creep at high-temperature, for example, the strain value of the as-cast alloy crept up to fracture is only about 7%, as shown in as shown in Fig. 4, the bigger plastic deformation occurs in the boundary regions with single γ phase, as shown in Fig. 7(b). This indicates that, compared to the lamellar α2/γ phases, a lower strength appears in the irregular serrated boundary regions with single γ phase in the alloy. Therefore, the strain of the alloy during creep occurs mainly in the irregular serrated boundary regions with single γ phase. Namely, the irregular serrated boundaries with single γ phase are thought to be the weak region of the creep strength during creep of the ascast high Nb-TiAl alloy. After heat treatment, the irregular serrated boundaries with single γ phase in the alloy disappear, and the grain boundary displays the smooth and straight features, as shown in Fig. 1(b), which may improve the creep resistance of alloy. Therefore, the creep life of the alloy after heat treatment increases from 147 h to 297 h under the applied stress of 200 MPa at 800 1C, as shown in Fig. 4, the increasing extent of the creep life is about 102%, which indicates that eliminating irregular serrated boundaries with single γ phase may improve the creep properties of the alloy to a great extent. The microstructure of both as cast and heat treated high NbTiAl alloy consists of lamellar α2/γ phases, and the lamellar α2/γ phases within grains display the wavy-like feature, as shown in Fig. 1(a) and (b). But after crept up to fracture, the lamellar α2/γ phases with wavy-like feature in alloys have transformed into the regular and straight configuration, as shown in Figs. 7(a) and 8, which indicates that the configuration transformation of the lamellar α2/γ phases occurs during creep at high temperatures due to the diffusion of the elements. 4.3. Analysis on initiation and propagation of crack during creep Although the irregular serrated boundaries with single γ phase may be eliminated, by means of heat treatment at high temperature, to enhance the creep properties of the alloy to a great extent, the crack in the alloy during creep is firstly initiated and propagated along the boundary, which suggests that the boundaries is still the weaker region of the creep strength in the alloy during creep. As creep goes on, significant amount of dislocations are activated in the alloy for piling up in the region near the boundaries, as shown in Fig. 9(b), which may cause the stress concentration to promote the initiation and propagation of the crack along the boundary up to the occurrence of creep fracture. Actually, the initiation of the crack may releases the stress concentration in the region. But the stress concentration may occur again in the region near the crack as creep goes on, which results in the propagation of the crack along the boundary. Especially, the grain boundaries perpendicular or at 451 angles relative to the stress axis bear the bigger shearing stress during creep, which causes easily the initiation and propagation of the
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cracks along the boundaries, as shown in Figs. 10 and 11. Although the boundaries with various orientations may stop the propagation of the crack when the one propagates to the triple junction region of the boundaries, the stress concentration occurs again in the region as the creep goes on, which may promote the initiation and propagation of the cracks occurring in the adjacent boundaries, as shown in Fig. 10(b) and (c). As the propagation of the multi-cracks in different cross-sections goes on, the tearing edges and secondary cracks are formed at the tip region of the primary crack along the direction with bigger shearing stress, which results in the gradual connection of the multi-cracks in different crosssections, by the tearing edges and secondary cleavage plane, up to the occurrence of creep fracture [36]. This is thought to be the fracture mechanism of the alloy during creep.
5. Conclusions (1) The microstructure of the as-cast high Nb-TiAl alloy consists of the lamellar γ⧸α2 phases with various orientations, and the irregular serrated boundaries with single γ phase are located in between the lamellar ones. After heat treatment, the irregular serrated single γ phase disappears in the boundary regions, and the regular boundaries are formed in between the lamellar α2/γ phases with various orientations. (2) Under the applied stress of 200 MPa at 800 1C, the creep life of as-cast high Nb-TiAl alloy is measured to be 145 h, and the one of the alloy after heat treated is enhanced to 297 h due to eliminating the irregular serrated boundaries with single γ phase. The creep activation energy of the heat treated alloy is measured to be Q¼ 432 kJ/mol in the ranges of the applied temperatures and stresses. (3) The deformation mechanism of alloy during creep is dislocations slipping in the matrix, the dislocations slipping in the matrix may be decomposed to form the configuration of the partials plus the stacking faults. The interaction of the deformation dislocations and stacking faults may hinder dislocations movement to improve the creep resistance of alloy. (4) The strain of as-cast alloy during creep occurs mainly in the irregular serrated boundary regions with single γ phase. After the alloy is heat treated, the primary/secondary slipping systems of dislocations are alternately activated in the latter stage of creep, which may result in the bigger plastic deformation for contorting the lamellar α2/γ phases in alloy. (5) The cracks in the alloys during creep are firstly initiated and propagated in the boundaries, which are parallel to the lamellar γ⧸α phases, along the direction perpendicular or at about 451 angles relative to the stress axis up to the occurrence of creep fracture, which is thought to be fracture mechanism of the alloy during creep.
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