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Creep behavior and deformation feature of TiAl–Nb alloy with various states at high temperature
Author’s Accepted Manuscript Creep Behavior and Deformation Feature of TiAlNb Alloy with Various States at High Temperature Sugui Tian, LV Xiaoxia, YU Huichen, Wang Qi, Jiao Zehui, Sun Haofang www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(15)30557-8 http://dx.doi.org/10.1016/j.msea.2015.10.096 MSA32947
To appear in: Materials Science & Engineering A Received date: 31 March 2015 Revised date: 20 October 2015 Accepted date: 25 October 2015 Cite this article as: Sugui Tian, LV Xiaoxia, YU Huichen, Wang Qi, Jiao Zehui and Sun Haofang, Creep Behavior and Deformation Feature of TiAl-Nb Alloy with Various States at High Temperature, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.10.096 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Creep Behavior and Deformation Feature of TiAl-Nb Alloy with Various States at High Temperature TIAN Sugui 1,a, LV Xiaoxia 1, YU Huichen 2, WANG Qi 1, JIAO Zehui 2, SUN Haofang 1 1. School of Materials Science and Engineering, Shenyang University of Technology, Shenyang 110870; China 2. 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 ABSTRACT: By means of creep properties measurement and microstructure observation, an investigation has been made into the creep behaviors and deformation features of TiAl-Nb alloy with various states at high temperature. Results show that the microstructure of as-cast alloys consists mainly of lamellar 2 phases with various orientations, the boundaries located in between the ones consist of single phase. The bigger plastic deformation of alloy during forging may refine the grains and increase the volume fraction of boundaries which consists of the fine block-like 2 phases with weaker strength. During creep, the bigger volume fraction of boundaries increase the probability of crack initiated and propagated along boundaries, which is thought to be the main reason of the forged alloy displaying a lower creep resistance. Compared to the forged alloy, the as-cast alloy displays a better creep resistance and longer creep life at high temperature. The deformation mechanism of as cast alloy during creep is significant amount of dislocations shearing into the lamellar phases in the form of dislocations rows and rings. Compared with Ti3Al phase, the phase possesses a weaker strength, therefore, the crack is easily initiated and propagated along the inclined boundaries. Wherein, the propagation of crack along the direction parallel to the lamellar phases displays the smooth surface of fracture, while the tearing edges appear on the surface of fracture in one side of the inclined boundaries relative to lamellar phases, which is attributed to the Ti3Al phase with better strength hindering the propagation of crack during creep. KEYWORDS: TiAl-Nb alloy, Microstructure, Creep, deformation mechanism, Fracture feature. 1. Introduction Because the TiAl alloys have an excellent strength, creep resistance and oxidation resistance at high temperatures [1-3], they are regarded as the better high-temperature structural materials with potential applications prospect [4-6], and are expected to make the hot parts in aero-engines for replacing the high-density metal materials [7,8]. But the fact that the TiAl alloys display a poor ductility at room temperature restricts their widely application. The plasticity of alloy at lower temperature is related to the chemical composition and solidification process of them. Some investigations indicated [9,10] that the solidification process of TiAl alloys includes a peritectic reaction (L + → The phases transformation processes of the alloy during solidification and cooling are described as follows : L → L + → → → → here, the phase is formed in the initial period of the solidification. Adding W element may improve the stability of the initial phase to favor the solidification of TiAl alloy [11,12]. Especially, adding small amount element W may improve the strength of TiAl alloys at room and high temperatures, and there is a little influence on plasticity of the one at room temperature [13], due to enhancing the transformation temperature of the toughness/brittleness.
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Moreover, the Ti atoms in alloy may react with trace B to form the fine TiB phase, which acts as the particles of the neterogeny nucleation to promote the uniform nucleation of phase and refine the grains of TiAl alloy [14]. And adding trace B may decrease the volume fraction of phase for enhancing the creep resistance of alloys. In particular, the creep resistance of alloy may be obviously improved when the volume fraction of phase is less than 12% [15]. In addition, adding trace Y element may both enhance the resistance oxidation and refine the grains of as-cast TiAl alloys [16]. Compared to the ordinary TiAl alloys, TiAl-Nb alloy has better strength, creep resistance and oxidation resistance at high temperature, therefore, the one is thought to be the excellent high-temperature structural materials for applying in aeronautics and astronautics field [17]. Actually it is thought that the creep resistance of alloys is related to their microstructure and deformation mechanism which includes the dislocations slipping and twinning deformation during creep [18]. Although the effects of casting technologies and alloying on microstructure and mechanical properties of TiAl alloys had been reported [19,20], the alloys with various microstructures display different creep resistance due to their deformation mechanism. Especially, the grain sizes of the as-cast TiAl-Nb alloy may be refined by means of the isothermal forged technology, but the creep behaviors and deformation mechanism of as-cast and forged TiAl-Nb alloys at high temperature is still unclear. In the paper, the TiAl-Nb based alloy are alloying treated by adding the traces W, B, and Y, and then some ingots of the alloy is isothermally forged to refine the grain size. Combined with creep properties measurement and microstructure observation, the creep behaviors and deformation mechanisms of the TiAl-Nb alloy with various states are investigated, which may provide the theory basis for promoting the development and application of TiAl-Nb alloys. 2. Experimental procedure The TiAl-Nb 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 parts of the ingot is isothermally forged into the square billet at 1250oC, the deformation quantity of the billet is about 80%, which is defined as the forged alloy. The nominal composition of the alloy is Ti-44Al-8Nb0.2W-0.2B-0.1Y. Then the ingot of the alloy is cut into some billets with the sizes of 14 mm 40 mm 40 mm, the billet of the as-cast and forged TiAl-Nb alloys were cut into the specimens with cross-section of 4.5 mm 2.5 mm and gauge length of 20 mm. And the uni-axial tensile creep tests were performed under constant load, in a GWT504-model creep testing machine, for measuring creep curves of the alloy under the applied various stresses at 900 ºC. Furthermore, the apparent creep active energy and apparent stress exponent of the alloy during steady state creep are calculated according to the data of creep curves. And then the microstructures of as-cast, forged and creep ruptured alloy are observed under scanning electron microscopy (SEM). Samples for transmission electron microscope (TEM) observation were cut into the discs along the direction plane parallel to the width plane of the creep samples and thinned down to ~50m mechanically. TEM foils were electro-polished in solution containing 5ml HF + 10ml HNO3 + 85ml H2O at -25oC. Wherein, TEM specimens of the as cast alloy were chosen in the regions about 3 mm away from the necked region of the creep samples to avoid the influence of the necking on dislocations configuration. The transmission electron microscope (TEM) with TECNA20 model is used to observe the microstructure evolution and deformation features of the alloy during creep at high temperature, and deformation and fracture mechanism of the alloy during creep are briefly discussed. 3. Experimental results and analysis 3.1 Microstructure of alloy The macro-morphology of the as-cast TiAl-Nb alloy is shown in Fig. 1(a), three grains in the picture 2
are denoted in the regions A, B and C, respectively, the sizes of the grains are about 500-800 m, the grains in alloy consists of lamellar structure with black and white phases. The lamellar phases within one grain are arranged along the same direction, but the lamellar phases within different grains are arranged along various orientations. The dark regions between the lamellar phases with different orientations are defined as the boundaries of alloy. The grain boundaries consisting of single phase [21] display the irregular serrated configuration, as marked in the region D.
The magnified morphology in another region is shown in Fig. 1(b), which consists of three grain marked by the letters E, F and G, the boundaries between the grains consist still of single phase with the irregular configuration. Some white needle-like rod precipitates are distributed in boundaries and within the grains, as marked by the arrows in Fig. 1(b). It is understood from Fig. 1(b) that the smaller spacing between the lamellar phases appears in the grains E and F, the bigger spacing between the lamellar ones displays in the grain G. Fig. 1 The TEM images of TiAl-Nb alloy is shown in Fig. 2, indicating that the microstructure of alloy consists of the lamellar double phases, the SAD pattern of the ones in the alloy, as marked by the letters D and E, is indicated in up-right region of the photo. According to the SAD pattern, the lamellar D phase is identified as -TiAl with Ll0 tetragonal structure, and the lamellar E phase is identified as α2-Ti3Al with D019 structure. The size of the lamellar phases in thickness is about 50-100 nm, the lamellar phases with parallel feature are alternately arranged along the same orientation. And it is thought that the coherent interface keeps in between the -TiAl and α2-Ti3Al phases due to no interfacial dislocations appearing in between the ones, which is well agreement with the result of the literatures [22, 23]. Fig. 2 The morphology of the forged TiAl-Nb alloy is shown in Fig. 3. It is indicated, compared with as cast alloy, the smaller grain size displays in the forged alloy, the sizes of the alloy are measured to be 80-120 m. Although the smooth and level lamellar phases still displays in the region F, the lamellar phases in some grains displays the distorted configuration, as shown in the regions D and E. Especially, the microstructure in the boundary regions displays the irregular and disordered configuration, the size of the boundaries in width is about 10m, The magnified morphology in the square frame of the grain boundary is shown in the down-right of Fig. 3, in which the lamellar -TiAl and α2-Ti3Al phases are transformed into the smaller block-like configuration. All morphology features stated above, including the smaller grains size, distorted lamellar configuration and boundaries, are attributed to the bigger plastic deformation of the alloy during isothermal forging. Fig. 3 3.2 Creep behaviors of alloy at high temperature Under the applied stress of 120MPa at 900 oC, the creep curves of the TiAl-Nb alloys with different states are measured as shown in Fig. 4, wherein, the creep curves of the as cast and forged alloys are marked by the numbers l and 2, respectively. Fig. 4 It is indicated from Fig. 4 that the smaller strain and bigger strain rate of the alloy occurs in the initial period of creep, and the strain rate of alloy decreases as the creep goes on. After crept for 10 h, the creep of alloy enters steady-state stage, the strain rate of alloy during steady-state creep is measured to be 0.0130%/h, the lasting time of alloy during steady state creep is about 150 h, the creep lifetime of alloy is 3
measured to be 237 h. Compared to the as cast alloy, the forged alloy during steady state creep displays a higher strain rate, the strain rate of the alloy during steady-state creep is measured to be 0.0329%/h, the lasting time of alloy during steady state creep is about 40 h, the creep life of alloy is measured to be 75h. Therefore, it may be concluded that, compared to the forged alloy, the as cast alloy displays a better creep resistance and longer creep life. The creep curves of the as-cast TiAl-Nb alloy at different conditions are measured, as shown in Fig. 5. The creep curves of alloy under the applied different stresses at 900 ºC are shown in Fig. 5(a), indicating that, the creep lifetime of alloy under the applied stress of 120 MPa is measured to be 237 h. The initial strain of the alloy increases slightly when the applied stress enhances to 130 MPa at 900 ºC, after crept for 20 h, the creep of alloy enters steady-state stage, the strain rate and lasting time of alloy during steady-state creep are measured to be about 0.0273%/h and 100 h, the creep lifetime of alloy is measured to be 163 h., When the applied stress increases to 140 MPa, the lasting time of the alloy during steady-state creep further shortens to about 35 h, the strain rate of alloy during steady-state creep is measured to be 0.0529%/h, and the creep lifetime of alloy is measured to be 76 h. The creep curves of the alloy under the applied stress of 140 MPa at different temperature are shown in Fig. 5(b), which indicates that the smaller strain of alloy occurs in the initial period of creep at 890 ºC. After crept for 10 h, the creep of alloy enters steady-state stage in which the lasting time of alloy is about 70 h, the strain rate of alloy during steady-state creep is measured to be 0.0472%/h, and the creep lifetime is measured to be 109 h. When the creep temperature enhances to 900 ºC, the creep lifetime of alloy is measured to be 76 h. When the creep temperature enhances to 910 ºC, the strain rate of the alloy during steady state creep increases to 0.0748%/h, the creep lifetime is measured to be 53.5 h. Fig. 5 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 alloy as the creep goes on, which decreases the strain rate of alloy 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 Norton-Baily law given as follows:
ss A An exp(
Q ) RT
(1)
Where ss is the strain rate during steady-state creep, A is the constant related to material structure, 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 Fig. 5, the strain rate of the as-cast 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 (ln ss 1/ T ) and (ln ss ln a ) , as shown in Fig. 6(a) and (b), respectively. Therefore, in the ranges of the applied temperatures and stresses, the apparent creep activation energy and stress exponent of the alloy during steady-state creep are measured to be Q = 413.9 kJ /mol and n = 7.8, respectively. According to the stress exponent, it may be deduced that the dislocations slipping and shearing the lamellar phases dominates the strain rate of alloy during steady-state creep. Fig. 6 3.3 Deformation features of alloy during creep After crept for 237 h up to rupture under the applied stress of 120 MPa at 900 ºC, the microstructures of as-cast TiAl-Nb alloy at different regions are shown in Fig. 7, the direction of the applied stress is 4
marked by the arrow in Fig. 7(a), which indicates that the various deformation features display in different regions of the sample. The smooth and level lamellar phases keep still in the region A far from fracture, as shown in Fig. 7(b), here, a set lamellar congeries with same orientation is defined as a grain, the boundaries are located in between the lamellar /congeries with various orientations, the lamellar phases within various grains are arranged along different orientations. Fig. 7 The microstructure in the region B is shown in Fig. 7(c), which indicates that significant amount of congeries in the region keep still the smooth and level lamellar configuration, but the lamellar/phases in local region displays the contorted configuration, as marked by the arrow in Fig. 7(c). Here, the boundaries with irregular serrated configuration consist still of the single phase, as marked by the letter H, and the lamellar congeries with smaller size and various orientations exist in the interior of grain, as marked in the region J. The microstructure in the region C near fracture is shown in Fig. 7(d), which consists still of the regular lamellar/phases, the grain boundaries are located in between the lamellar / congeries with various orientations. Because the bigger plastic deformation occurs in the region, the original smooth and level lamellar / phases had been contorted to display the tortuous configuration, as marked by the arrow in Fig. 7(d), and the lamellar /phases in the region near boundaries display the disorder configuration, as marked in the region L. After crept up to fracture of the forged alloy at 900 oC/120MPa, the microstructure in the region near fracture are shown in Fig. 8. It is shown that the original lamellar phases in some regions still keep the smooth and level lamellar configuration, as marked by the letter K in Fig. 8(a), and the lamellar phases in different grains display the various orientations. If the lamellar congeries with various orientations are defined as different grains, the lamellar congeries in the grains display the various spacing distances. The regions between the lamellar congeries with different orientations are defined as grain boundaries which consist of fine block-like phases, as marked by the letter M in Fig. 8(a), and the size of boundaries in width increase to about 10-20m. Moreover, some needle-like precipitates are distributed in the lamellar phases and boundary regions, as marked by white arrow in Fig. 8(a). Fig. 8 Due to the bigger plastic deformation occurring in another region, the lamellar phases in some grains displays the distorted configuration, as shown in the regions N, P and Q of Fig. 8(b). Moreover, the fine block-like phases are distributed in the boundary regions, as marked by the thick arrow, and the crack appears in the boundary region which consists of fine block-like phases, as marked by the fine arrow, which may be thought to be the initiation of crack along boundaries. As the creep goes on, the crack is propagated along boundaries to form the bigger crack, as marked by the longer arrow in Fig. 8(c), and a smaller crack in the boundary region is marked by the fine arrow. When the bigger crack is further propagated along the boundary as the creep continues, the one may connect with the smaller crack to promote the propagation of crack up to the occurrence of creep fracture, which is thought to be the fracture mechanism of the forged alloy during creep. Therefore, it may be concluded that the boundaries in the forged alloy are still the weaker region of creep resistance during creep at high temperature. After crept for 237 h up to rupture under the applied stress of 120 MPa at 900 ºC, the microstructure of as-cast TiAl-Nb alloy are shown in Fig. 9, significant amount of dislocation rows appear in the local region of lamellar phases, the dislocation lines in the dislocation rows are parallel to each other, as marked by the black arrow in Fig. 9(a). The dislocation loop appears in the dislocation rows, as marked by the white arrow, the magnified morphology of the loop is indicated in the right-top of Fig. 9(a). Moreover, the regular dislocation rows distribute in the interface of the lamellar /phases, as shown in the square frame of Fig. 5
9(a), the magnified morphology of the dislocation rows is shown in the left down of Fig. 9(a). Moreover, some white granular particles are precipitated in the matrix, as marked by the short arrow in Fig. 9(a). In local region of sample, the lamellar phases are parallel to each other, as shown in Fig. 9(b), the size of the lamellar/phases is about 500—700m. Some dislocations have sheared into the lamellar phases, as marked by the black arrow in Fig. 9(b). Moreover, significant amount of dislocations are distributed in the interfaces of the lamellar phases to form the regular dislocations networks, as shown in the square frame, the magnified morphology in the square frame region is shown in the right down of Fig. 9(b). It is thought by analysis that the creep dislocations moving to the interface of phases may react with the dislocation network to change the original moving direction, which may promote the climbing of dislocations. Therefore, the dislocations networks located in the interfaces have the coordinating role on the deformation hardening and recovery softening of alloy during creep. Fig. 9 In another local region of sample, dislocation shearing into lamellar phases is marked by the arrow in Fig. 9(c), the dislocations located in the middle region of the photo display the dislocation rows configuration with parallel feature. Furthermore, the inclined boundary is located in between the lamellar phases, and the isoclinic stripes appearing in the inclined boundary are parallel to the orientation of lamellar phases. Some dislocations shear through the interface of lamellar phases, as shown in the square frame, and the magnified morphology in the square frame region is shown in the right top of Fig. 9(c). The fact that the creep dislocations within lamellar phases stop in the interface of phases indicates that the phase boundaries may hinder the slipping of dislocations. Therefore, it may be concluded that the deformation mechanism of the alloy during creep is dislocations slipping in the phases, wherein, the dislocations slipping in the matrix may react each other to form the dislocations networks, which may promote the climbing of dislocations. 3.4 Initiation and propagation of crack After as-cast TiAl-Nb alloy is crept for 237 h up to rupture at 900 ºC/120 MPa, the morphologies of the cracks initiating and propagating along boundaries near fracture are shown in Fig. 10, in which the direction of the applied stress is marked by the arrows. The boundary in the region is about 45 angles relative to the stress axis, as marked by straight line in Fig. 10(a), the orientation of the lamellar phases in the down side of the boundary is parallel to the direction of the boundary, but the orientation of the lamellar phases in the upper region of the boundary is perpendicular to the stress axis. Wherein, the cavity formed along the boundary of alloy during creep is marked by the fine arrow, which is regarded as the initiation of crack, the bigger crack along the boundary is regarded as the propagation of the crack, as marked by the bigger arrow in Fig. 10(a). It is understood from Fig. 10(a) that the propagating direction of the crack on one side is parallel to the orientation of the lamellar phases. In another region near fracture, the bigger crack along the boundary is still parallel to the orientation of the lamellar phases, and the boundary is at about 45 angles relative to the stress axis, as marked by the arrow in Fig. 10(b), but the propagating direction of the crack in another side is perpendicular to the orientation of the lamellar phases. This indicates that, during creep, the crack in the alloy is easily initiated and propagated along the boundaries at about 45 angles relative to the stress axis, which is attributed to the 45 angles boundaries bearing a bigger shear stress and its weaker strength. Fig. 10 6
In another region near fracture, the fracture surface originating from the crack propagation displays the various configurations. On the upper surface of the crack, the propagating direction of crack is inclined with the orientation of the lamellar phases, but on the down surface of the crack, the propagating direction of crack is parallel to the orientation of the lamellar phases, asshown in Fig. 10(c). The down surface of the crack formed during propagation displays the smooth feature, which suggests that the propagation of crack experiences a smaller resistance. But the fracture surface of displaying non smooth and tearing edges appears in the upper surface of the crack, as marked by the thicker arrow in Fig. 10(c), which suggests that the propagation of crack experiences a bigger resistance. It is thought by analysis that the microstructure of TiAl–Nb alloy consist of lamellar 2 phases, compared with Ti3Al, the strength of phase is weaker, therefore, the Ti3Al is the strengthening phase of the alloy. The grain boundary consisting of single phase is the weaker region of strength during creep of alloy, moreover, the boundaries at about 45 angles relative to the stress axis bear the bigger shear stress under loading. Therefore, the initiation and propagation of crack occurs easily in the boundaries along the direction at about 45 angles relative to the stress axis during creep. Furthermore, the resistance of the crack propagation is related to the orientation of the lamellar phases. When the direction of the crack propagation is parallel to the orientation of the lamellar phases, the propagation of crack along the single phase with weaker strength encounters only a smaller resistance. Therefore, the fracture of alloy after crept up to rupture displays a smooth surface, as marked by the fine arrow in Fig. 10(c). When the direction of the crack propagation is inclined with the orientation of the lamellar phases, the propagation of crack may encounter the Ti3Al phase with excellent strength to increase the resistance of crack propagation during creep, which results in the tearing edges appearing in the surface of fracture, as marked by the thicker arrow in Fig. 10(c). This is well consistent with the experimental results above state. 4. Discussion 4.1 Theory analysis of creep resistance of alloy at various states The microstructure of as-cast TiAl-Nb alloy consists of lamellar phases with different orientations, as shown in Fig. 1, wherein, the grain boundaries of alloy consists of single phase with smaller size in thickness. Compared to phase, the phase possesses a better strength, therefore, the deformation of the alloy during creep occurs mainly in phase. And it is thought that the characteristic of phases being easily deformed may delay the stress concentration resulting from the inharmonious deformation of the lamellar phases during creep. During isothermal forging, the bigger plastic deformation occurs in the alloy, which diminishes the grain size of them from 500m to about 80-120m. Therefore, compared to as cast TiAl-Nb alloy, the alloy after isothermal forged increases the quantity and volume fraction of the grain boundaries to a great extent. Especially, the wider boundary of the forged alloy consists of fine block-like phases with irregular features. During creep, the resistance of dislocations slipping in the boundary regions is smaller, the stress concentration originating from the large number dislocations piled up in the regions near boundary may results in the initiation and propagation of cracks along the boundaries, as shown in Fig. 8(b) and (c). Therefore, the wider boundaries of alloy are still the weaker regions of strength during creep at high temperature. It is thought by analysis that on the one hand, the bigger plastic deformation during isothermal forged refines the grains size of alloy, on the other hand, the process makes the original smooth and level lamellar phases transforming into the distorted configuration. Namely, the deformation of alloy during forging results in both decreasing creep resistance of microstructure and increasing the volume fraction of boundaries with weaker strength, so that the probability of the crack initiating and propagating along the
7
boundaries during creep at high temperature increases, which is thought to be the main reason of the forged alloy displaying a lower creep resistance and shorter lifetime. 4.2 Microstructure evolution and deformation features during creep The microstructure of as-cast TiAl-Nb alloy consists of lamellar /2 phases with different orientations, the size of lamellar phases in thickness is about 50-100 nm, as shown in Fig. 2, and the lamellar phases keep the coherent interface due to no interfacial dislocations distributed in between them. But crept for 237h up to rupture under the applied stress of 120 MPa at 900 ºC, the size of the lamellar phases in thickness increases to about 500-800 nm, as shown in Fig. 9(b), which suggests that the size of the lamellar phases in thickness increases as the creep goes on. Although the lamellar phases before creep keep the coherent interface, the interfacial dislocations have appeared in between the lamellar phases after crept for 237 h at high temperature, as marked by the square frame in Fig. 9(a) and (b). This indicates that the coherent interface in the lamellar phases has transformed into the semi-coherent interface as the size of the lamellar phases increases. It is thought by analysis that compared to the coherent interface, the extent of the lattice strain in the interface of phases decreases after the one is transformed into the semi-coherent interface, therefore, the resistance of dislocations in one phase shearing into another phase diminishes. Compared with Ti3Al phase, the phase possesses a weaker strength [22], so the plastic deformation of alloy during creep occurs firstly in the phase with weaker strength. Significant amount of dislocations may firstly activate in phase, a few dislocations appears in phase, as shown in Fig. 9(a) and (b). And the quantity of dislocations shearing into the lamellar phase increases with the strain of alloy, wherein, the dislocations both shear into the phase in the form of the rows and slip in the phase in the form of ring, as marked by the arrow in Fig. 9(a) and (b). As the creep goes on, significant amount of dislocations slipping to the interface are hindered to pile up in the region near the interface, which may bring the stress concentration. The dislocations piled up in phase may shear into the phase along the interface dislocation network when the value of the stress concentration exceeds the strength of 2 phase. Moreover, significant amount of dislocations shearing into the lamellar phases are alternately activated, which may twist the lamellar phases within grains, as marked by the white arrow in Fig. 7(c), to increase the strain of alloy. Moreover, the grain boundaries in between the lamellar congeries display the irregular serrated configuration, and the ones consist of the single phase with Ll0 tetragonal structure [21], as marked by the arrow in Fig. 1(a). Because the TiAl phase possesses a weaker strength and better ductility, the bigger plastic deformation occurs in the boundary region of as-cast alloy during creep at high temperature. The boundary regions with wider feature are marked in the regions H and J of Fig 7(c), which suggests that the grain boundary is a weaker link of restricting the creep strength of alloy. It is indicated from Fig. 8 and 10 that the initiation and propagation of crack occurs mainly in the grain boundaries at about 45 angles relative to the stress axis during creep of alloy due to the direction bearing a bigger shear stress. It is thought by analysis that the boundaries of as cast and forged alloys consist of single phase and fine block-like phases with weaker strength, during creep at high temperature, the bigger shear stress may promote the initiation of crack along the boundaries with weaker strength. And as the creep goes on, the bigger shear stress promotes the propagation of crack along the boundaries up to the occurrence of creep fracture. Fig. 11 The schematic diagram of crack initiated and propagated along the boundary at 45 angles relative to the stress axis is shown in Fig. 11, the direction of the applied stress is marked by the arrows. Because the boundaries of alloys possess a weaker strength, the initiation of the crack occurs firstly in the inclined 8
boundary under the action of the bigger shearing stress, which is schematically shown in Fig. 11(a). As the creep goes on, the size of the crack propagating along boundary increases under the action of bigger shearing stress, as shown in Fig. 11(b). And then the crack is further propagated along the inclined boundary until encountering the boundary with another orientation, as marked by the arrow in Fig. 10(b), which may hinder the propagation of the crack, as schematically shown in Fig. 11(c). Although the grain boundary is the weaker link of strength during creep of the as cast alloy, the propagation of the cracks along boundary may encounter the Ti3Al phase with better strength to hinder the one of them. Therefore, when the orientation of the lamellar phases is inclined with the direction of the grain boundaries, the inclined lamellar Ti3Al phase relative to the boundary may restrain the propagation of the cracks to improve the creep resistance and life of as cast alloy. The analysis stated above is well agreement with the experiment results. 4. Conclusion (1) The microstructure of as-cast TiAl-Nb alloys consists mainly of lamellar 2 phases, the boundaries with irregular serrated configuration consist of single phase, and the ones are located in between the lamellar 2 phases with various orientations. (2) The bigger plastic deformation of alloy during isothermal forging may refine the size of the grains and increase the volume fraction of boundaries which consists of the fine block-like 2 phases with weaker strength. During creep of the forged alloy, the wider size and bigger volume fraction of boundaries increase the probability of crack initiated and propagated along boundaries, which is thought to be the main reason of the forged alloy displaying a lower creep resistance and shorter lifetime. (3) During creep at temperature near 900 oC, compared to forged alloy, the as-cast alloy displays a better creep resistance and longer creep life. The deformation mechanism of alloys during creep is significant amount of dislocations shearing into the lamellar phases in the form of dislocations rows and rings. Moreover, the dislocations networks located in the interfaces may promote the climbing of dislocations to retard the stress concentration and improve the creep resistance of alloy. (4) Compared with Ti3Al phase, the phase possesses a weaker strength, therefore, the crack is easily initiated and propagated along the boundaries at about 45 angles relative to the stress axis until the occurrence of creep fracture. Wherein, the fracture parallel to the lamellar phases displays the smooth surface, while the tearing edges appear on the surface of fracture in one side of the inclined boundaries relative to lamellar phases, which is attributed to the Ti3Al phase with better strength hindering the propagation of crack during creep. Acknowledgements Sponsorship of this research by the National Basic Research Program of China under Grant No. 2011CB605506 is gratefully acknowledged. References [1] Zheng Z, Jiang Y, Dong HX, Tang LM, He YH, Huang BY. Environmental corrosion resistance of intermetallic compounds. Trans. Nonferrous Met Soc China 2009; 19: 582-585. [2] Ye XC, Su YQ, Guo JJ, Zhang LH, Hao ZJ. Effect of heat treatment on microstructure of TiAl based alloy casting by suction casting process, Trans Mater Heat Treat 2012; 33: 132-135. [3] Lu X, Zhao LH, Zhu LP, Zhang B, Qu XH. High-temperature mechanical properties and deformation behavior of high Nb containing TiAl alloys fabricated by spark plasma sintering. Mater 2012; 19(4): 354-359. 9
[4] Ding XF, Lin JP, Zhang LQ, Su YQ, Chen GL. Microstructural control of TiAl–Nb alloys by directional solidification. Acta Mater 2012; 60(2): 498-506. [5] Hénaff G, Gloanec AL. Fatigue properties of TiAl alloys. Intermetallics 2005; 13: 543-558. [6] Lin JP, Xu XJ, Wang YL, He SF, Zhang Y, Song XP, Chen GL. High temperature deformation behaviors of a high Nb containing TiAl alloy. Intermetallics 2007; 15: 668–674. [7] Kong FT, Chen YY, Tian J, Chen ZY. Development in TiAl Based intermetallics research. Master Sci Technol 2003; 11: 441-444. [8] Chen GQ, Zhang BG, Liu W, Feng JC. Crack formation and control upon the electron beam welding of TiAl-based alloys. Intermetallics 2011; 19: 1857-1863. [9] Kelly TJ, Juhas MC, Huang SC. Effect of a B2/gamma structure on the tensile properties of the cast gamma titanium aluminide Ti-48Al-2Cr-2Nb. Scr Metall Mater 1993; 29: 1409-1414. [10] McCullough C, Valencia JJ, Levi CG, Mehrabian R. Phase equilibria and solidification in Ti-Al alloys. Acta Metall Master 1989; 37: 1321-1336. [11] Naka S, Thomas M, Sanchez C, Khan, T. Development of third generation role of solidification paths. TMS 1997; 313-322. [12] Sun FS, Cao CX, Yan MQ, Kim SE, Yang TL. Alloying mechanism of beta stabilizers in a-TiAl alloy. Metall Mater Trans A 2001; 32: 1573-1589. [13] Liu ZC, Lin PJ, Chen GL. Effect of the addition W on the microstructure and mechanical properties for high-Nb TiAl alloy. Trans Mater Heat Treat 2001; 22: 8-13. [14] Hu D, Yang C, Huang A, Dixon M, Hecht U. Grain refinement in beta-solidifying Ti44A18Nb1B. Intermetallics 2012; 23: 49-56. [15] Lin JG, Zhang YG, Chen CQ. Effects of the microstructure of γ-TiAl alloys on the creep behavior. Journal of Beijing University of Aeronautics and Astronautics 1998; 24: 734-737. [16] Chen YY, Kong FT, Han JC, Chen ZY, Tian J. Influence of yttrium on microstructure, mechanical properties and deformability of Ti-43A1-9V alloy. Intermetallics 2005; 13: 263-266. [17] Yan YQ, Zhou L, Wang WS, Zhang YN. 8.5 Nb–TiAl alloy with fine grains. Journal of alloys and compounds 2003; 361: 241-246. [18] Wang YH, Lin JP, He YH, Wang YL, Chen GL. Diffusion behavior of Nb element in high NB containing TiAl alloys by reactive hot pressing. Rare Met 2006; 25: 349-354. [19] Kim HY, Maruyama K. Changes in lamellar microstructure by parallel twinning during creep in soft PST crystal of TiAl alloy. Trans Nonferrous Met Soc China 2002; 12: 561-568. [20] Niu HZ, Chen YY, Xiao SL, Xu J. Microstructure and mechanical poperties of a novel beta g-TiAl alloy. Intermetallics 2012; 31: 225-231. [21] Qu HP, Li P, Zhang SQ, Li A, Wang HM. The effects of heat treatment on the microstructure and mechanical property of laser melting deposition γ-TiAl intermetallic alloys. Materials and Design 2010; 31: 2201-2210. [22] Li XW, Sun HF, Zhang P, Fang WB. The effect of strain on dynamic recrystallization of PM Ti–45Al10Nb intermetallics during isothermal forging. Intermetallics 2014; 55: 90-94. [23] Tian SG, Wang Q, Yu HC, Sun HF, Li QY. Microstructure and creep behaviors of a high Nb-TiAl intermetallic compound based alloy. Mater. Sci Eng. A 614, 2014: 338-346.
10
(a)
(b)
40m
20m
Fig. 1 Microstructure of the as-cast TiAl-Nb alloy. (a) macro-morphology, (b) magnified morphology.
Fig. 2 After heat treatment, microstructure of TiAl-Nb alloy.
11
Fig. 3 Microstructure of the isothermal forged TiAl-Nb alloy.
Strain,ε(%)
20 15
1 -- As-cast alloy 2 -- Forged alloy 120 MPa o T -- 900 C
10 5
0
2
1
50
100 150 200 250 Time, (h) Fig. 4 Creep curves of alloy with various states under the applied stress of 120MPa at 900 ℃
12
30
30
18
24 Strain,
Strain,ε(%)
24
-- 140 MPa
1 -- 120 MPa 2 -- 130 MPa 3 -- 140 MPa o T -- 900 C
12 3
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1 -- 890 oC o 2 -- 900 C o 3 -- 910 C
12 3
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2
6
1
(b)
(a) 0
50
100 150 Time, (h)
200
0
250
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40
60 80 Time, (h)
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Fig. 5 Creep curves of alloy at different conditions. (a) Applied different stresses at 900 ºC, (b) applied stress of 140MPa at different temperatures.
-3.30
-3.3
Q = 413.9 kJ/mol
-3.45
-3.6
ln ss
ln ss
-3.60 -3.75 -3.90 -4.05 -4.20
n = 7.80
-3.9 -4.2 -4.5
(a) 8.46
(b) 8.49
8.52 8.55 1/T,10-3K-1
8.58
8.61
-4.8
4.80
4.84
4.88
4.92
4.96
In
Fig. 6 Dependence of the strain rates of alloy during steady state creep on temperatures and stresses. (a) Strain rates and temperatures, (b) strain rates and applied stress
13
Fig. 7 Microstructure in different regions of sample after crept up to fracture at 120 MPa/900 ºC. (a) Schematic diagram of marking observed regions in sample, (b), (c), (d) being SEM morphologies corresponding to A, B, C regions, respectively.
Fig. 8 Microstructure in different regions of the forged sample after crept up to fracture at 120 MPa/900 ºC. (a) lamellar congeries in different grains, (b) crack initiated along boundaries, (c) the propagation of crack.
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Fig. 9 Microstructure of alloy crept for 237h up to fracture at 900 ºC/120 MPa. (a) dislocation rows, (b) interfacial dislocations networks, (c) dislocations shearing lamellar phase.
Fig. 10 After crept for 237 h up to fracture at 900oC/120MP, initiation and propagation of cracks. (a) Initiation of crack along the boundary at about 45 angles relative to the stress axis, (b) propagation of crack along the boundary, (c) non smooth surface formed in another side of crack
15
Fig. 11 Schematic diagram of crack initiated and propagated along boundary at 45 angle relative to the stress axis during creep at 900 oC. (a) Initiation of crack along boundary, (b) propagation of crack along boundary, (c) propagation of crack being stopped in the boundary with another orientation,
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PII: DOI: Reference:
S0921-5093(15)30557-8 http://dx.doi.org/10.1016/j.msea.2015.10.096 MSA32947
To appear in: Materials Science & Engineering A Received date: 31 March 2015 Revised date: 20 October 2015 Accepted date: 25 October 2015 Cite this article as: Sugui Tian, LV Xiaoxia, YU Huichen, Wang Qi, Jiao Zehui and Sun Haofang, Creep Behavior and Deformation Feature of TiAl-Nb Alloy with Various States at High Temperature, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.10.096 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Creep Behavior and Deformation Feature of TiAl-Nb Alloy with Various States at High Temperature TIAN Sugui 1,a, LV Xiaoxia 1, YU Huichen 2, WANG Qi 1, JIAO Zehui 2, SUN Haofang 1 1. School of Materials Science and Engineering, Shenyang University of Technology, Shenyang 110870; China 2. 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 ABSTRACT: By means of creep properties measurement and microstructure observation, an investigation has been made into the creep behaviors and deformation features of TiAl-Nb alloy with various states at high temperature. Results show that the microstructure of as-cast alloys consists mainly of lamellar 2 phases with various orientations, the boundaries located in between the ones consist of single phase. The bigger plastic deformation of alloy during forging may refine the grains and increase the volume fraction of boundaries which consists of the fine block-like 2 phases with weaker strength. During creep, the bigger volume fraction of boundaries increase the probability of crack initiated and propagated along boundaries, which is thought to be the main reason of the forged alloy displaying a lower creep resistance. Compared to the forged alloy, the as-cast alloy displays a better creep resistance and longer creep life at high temperature. The deformation mechanism of as cast alloy during creep is significant amount of dislocations shearing into the lamellar phases in the form of dislocations rows and rings. Compared with Ti3Al phase, the phase possesses a weaker strength, therefore, the crack is easily initiated and propagated along the inclined boundaries. Wherein, the propagation of crack along the direction parallel to the lamellar phases displays the smooth surface of fracture, while the tearing edges appear on the surface of fracture in one side of the inclined boundaries relative to lamellar phases, which is attributed to the Ti3Al phase with better strength hindering the propagation of crack during creep. KEYWORDS: TiAl-Nb alloy, Microstructure, Creep, deformation mechanism, Fracture feature. 1. Introduction Because the TiAl alloys have an excellent strength, creep resistance and oxidation resistance at high temperatures [1-3], they are regarded as the better high-temperature structural materials with potential applications prospect [4-6], and are expected to make the hot parts in aero-engines for replacing the high-density metal materials [7,8]. But the fact that the TiAl alloys display a poor ductility at room temperature restricts their widely application. The plasticity of alloy at lower temperature is related to the chemical composition and solidification process of them. Some investigations indicated [9,10] that the solidification process of TiAl alloys includes a peritectic reaction (L + → The phases transformation processes of the alloy during solidification and cooling are described as follows : L → L + → → → → here, the phase is formed in the initial period of the solidification. Adding W element may improve the stability of the initial phase to favor the solidification of TiAl alloy [11,12]. Especially, adding small amount element W may improve the strength of TiAl alloys at room and high temperatures, and there is a little influence on plasticity of the one at room temperature [13], due to enhancing the transformation temperature of the toughness/brittleness.
Corresponding author. +86-24-25494089(office), +86-13889121677 (Mobile); fax: +86-24- 25496768; E-mail address: [email protected] 1
Moreover, the Ti atoms in alloy may react with trace B to form the fine TiB phase, which acts as the particles of the neterogeny nucleation to promote the uniform nucleation of phase and refine the grains of TiAl alloy [14]. And adding trace B may decrease the volume fraction of phase for enhancing the creep resistance of alloys. In particular, the creep resistance of alloy may be obviously improved when the volume fraction of phase is less than 12% [15]. In addition, adding trace Y element may both enhance the resistance oxidation and refine the grains of as-cast TiAl alloys [16]. Compared to the ordinary TiAl alloys, TiAl-Nb alloy has better strength, creep resistance and oxidation resistance at high temperature, therefore, the one is thought to be the excellent high-temperature structural materials for applying in aeronautics and astronautics field [17]. Actually it is thought that the creep resistance of alloys is related to their microstructure and deformation mechanism which includes the dislocations slipping and twinning deformation during creep [18]. Although the effects of casting technologies and alloying on microstructure and mechanical properties of TiAl alloys had been reported [19,20], the alloys with various microstructures display different creep resistance due to their deformation mechanism. Especially, the grain sizes of the as-cast TiAl-Nb alloy may be refined by means of the isothermal forged technology, but the creep behaviors and deformation mechanism of as-cast and forged TiAl-Nb alloys at high temperature is still unclear. In the paper, the TiAl-Nb based alloy are alloying treated by adding the traces W, B, and Y, and then some ingots of the alloy is isothermally forged to refine the grain size. Combined with creep properties measurement and microstructure observation, the creep behaviors and deformation mechanisms of the TiAl-Nb alloy with various states are investigated, which may provide the theory basis for promoting the development and application of TiAl-Nb alloys. 2. Experimental procedure The TiAl-Nb 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 parts of the ingot is isothermally forged into the square billet at 1250oC, the deformation quantity of the billet is about 80%, which is defined as the forged alloy. The nominal composition of the alloy is Ti-44Al-8Nb0.2W-0.2B-0.1Y. Then the ingot of the alloy is cut into some billets with the sizes of 14 mm 40 mm 40 mm, the billet of the as-cast and forged TiAl-Nb alloys were cut into the specimens with cross-section of 4.5 mm 2.5 mm and gauge length of 20 mm. And the uni-axial tensile creep tests were performed under constant load, in a GWT504-model creep testing machine, for measuring creep curves of the alloy under the applied various stresses at 900 ºC. Furthermore, the apparent creep active energy and apparent stress exponent of the alloy during steady state creep are calculated according to the data of creep curves. And then the microstructures of as-cast, forged and creep ruptured alloy are observed under scanning electron microscopy (SEM). Samples for transmission electron microscope (TEM) observation were cut into the discs along the direction plane parallel to the width plane of the creep samples and thinned down to ~50m mechanically. TEM foils were electro-polished in solution containing 5ml HF + 10ml HNO3 + 85ml H2O at -25oC. Wherein, TEM specimens of the as cast alloy were chosen in the regions about 3 mm away from the necked region of the creep samples to avoid the influence of the necking on dislocations configuration. The transmission electron microscope (TEM) with TECNA20 model is used to observe the microstructure evolution and deformation features of the alloy during creep at high temperature, and deformation and fracture mechanism of the alloy during creep are briefly discussed. 3. Experimental results and analysis 3.1 Microstructure of alloy The macro-morphology of the as-cast TiAl-Nb alloy is shown in Fig. 1(a), three grains in the picture 2
are denoted in the regions A, B and C, respectively, the sizes of the grains are about 500-800 m, the grains in alloy consists of lamellar structure with black and white phases. The lamellar phases within one grain are arranged along the same direction, but the lamellar phases within different grains are arranged along various orientations. The dark regions between the lamellar phases with different orientations are defined as the boundaries of alloy. The grain boundaries consisting of single phase [21] display the irregular serrated configuration, as marked in the region D.
The magnified morphology in another region is shown in Fig. 1(b), which consists of three grain marked by the letters E, F and G, the boundaries between the grains consist still of single phase with the irregular configuration. Some white needle-like rod precipitates are distributed in boundaries and within the grains, as marked by the arrows in Fig. 1(b). It is understood from Fig. 1(b) that the smaller spacing between the lamellar phases appears in the grains E and F, the bigger spacing between the lamellar ones displays in the grain G. Fig. 1 The TEM images of TiAl-Nb alloy is shown in Fig. 2, indicating that the microstructure of alloy consists of the lamellar double phases, the SAD pattern of the ones in the alloy, as marked by the letters D and E, is indicated in up-right region of the photo. According to the SAD pattern, the lamellar D phase is identified as -TiAl with Ll0 tetragonal structure, and the lamellar E phase is identified as α2-Ti3Al with D019 structure. The size of the lamellar phases in thickness is about 50-100 nm, the lamellar phases with parallel feature are alternately arranged along the same orientation. And it is thought that the coherent interface keeps in between the -TiAl and α2-Ti3Al phases due to no interfacial dislocations appearing in between the ones, which is well agreement with the result of the literatures [22, 23]. Fig. 2 The morphology of the forged TiAl-Nb alloy is shown in Fig. 3. It is indicated, compared with as cast alloy, the smaller grain size displays in the forged alloy, the sizes of the alloy are measured to be 80-120 m. Although the smooth and level lamellar phases still displays in the region F, the lamellar phases in some grains displays the distorted configuration, as shown in the regions D and E. Especially, the microstructure in the boundary regions displays the irregular and disordered configuration, the size of the boundaries in width is about 10m, The magnified morphology in the square frame of the grain boundary is shown in the down-right of Fig. 3, in which the lamellar -TiAl and α2-Ti3Al phases are transformed into the smaller block-like configuration. All morphology features stated above, including the smaller grains size, distorted lamellar configuration and boundaries, are attributed to the bigger plastic deformation of the alloy during isothermal forging. Fig. 3 3.2 Creep behaviors of alloy at high temperature Under the applied stress of 120MPa at 900 oC, the creep curves of the TiAl-Nb alloys with different states are measured as shown in Fig. 4, wherein, the creep curves of the as cast and forged alloys are marked by the numbers l and 2, respectively. Fig. 4 It is indicated from Fig. 4 that the smaller strain and bigger strain rate of the alloy occurs in the initial period of creep, and the strain rate of alloy decreases as the creep goes on. After crept for 10 h, the creep of alloy enters steady-state stage, the strain rate of alloy during steady-state creep is measured to be 0.0130%/h, the lasting time of alloy during steady state creep is about 150 h, the creep lifetime of alloy is 3
measured to be 237 h. Compared to the as cast alloy, the forged alloy during steady state creep displays a higher strain rate, the strain rate of the alloy during steady-state creep is measured to be 0.0329%/h, the lasting time of alloy during steady state creep is about 40 h, the creep life of alloy is measured to be 75h. Therefore, it may be concluded that, compared to the forged alloy, the as cast alloy displays a better creep resistance and longer creep life. The creep curves of the as-cast TiAl-Nb alloy at different conditions are measured, as shown in Fig. 5. The creep curves of alloy under the applied different stresses at 900 ºC are shown in Fig. 5(a), indicating that, the creep lifetime of alloy under the applied stress of 120 MPa is measured to be 237 h. The initial strain of the alloy increases slightly when the applied stress enhances to 130 MPa at 900 ºC, after crept for 20 h, the creep of alloy enters steady-state stage, the strain rate and lasting time of alloy during steady-state creep are measured to be about 0.0273%/h and 100 h, the creep lifetime of alloy is measured to be 163 h., When the applied stress increases to 140 MPa, the lasting time of the alloy during steady-state creep further shortens to about 35 h, the strain rate of alloy during steady-state creep is measured to be 0.0529%/h, and the creep lifetime of alloy is measured to be 76 h. The creep curves of the alloy under the applied stress of 140 MPa at different temperature are shown in Fig. 5(b), which indicates that the smaller strain of alloy occurs in the initial period of creep at 890 ºC. After crept for 10 h, the creep of alloy enters steady-state stage in which the lasting time of alloy is about 70 h, the strain rate of alloy during steady-state creep is measured to be 0.0472%/h, and the creep lifetime is measured to be 109 h. When the creep temperature enhances to 900 ºC, the creep lifetime of alloy is measured to be 76 h. When the creep temperature enhances to 910 ºC, the strain rate of the alloy during steady state creep increases to 0.0748%/h, the creep lifetime is measured to be 53.5 h. Fig. 5 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 alloy as the creep goes on, which decreases the strain rate of alloy 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 Norton-Baily law given as follows:
ss A An exp(
Q ) RT
(1)
Where ss is the strain rate during steady-state creep, A is the constant related to material structure, 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 Fig. 5, the strain rate of the as-cast 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 (ln ss 1/ T ) and (ln ss ln a ) , as shown in Fig. 6(a) and (b), respectively. Therefore, in the ranges of the applied temperatures and stresses, the apparent creep activation energy and stress exponent of the alloy during steady-state creep are measured to be Q = 413.9 kJ /mol and n = 7.8, respectively. According to the stress exponent, it may be deduced that the dislocations slipping and shearing the lamellar phases dominates the strain rate of alloy during steady-state creep. Fig. 6 3.3 Deformation features of alloy during creep After crept for 237 h up to rupture under the applied stress of 120 MPa at 900 ºC, the microstructures of as-cast TiAl-Nb alloy at different regions are shown in Fig. 7, the direction of the applied stress is 4
marked by the arrow in Fig. 7(a), which indicates that the various deformation features display in different regions of the sample. The smooth and level lamellar phases keep still in the region A far from fracture, as shown in Fig. 7(b), here, a set lamellar congeries with same orientation is defined as a grain, the boundaries are located in between the lamellar /congeries with various orientations, the lamellar phases within various grains are arranged along different orientations. Fig. 7 The microstructure in the region B is shown in Fig. 7(c), which indicates that significant amount of congeries in the region keep still the smooth and level lamellar configuration, but the lamellar/phases in local region displays the contorted configuration, as marked by the arrow in Fig. 7(c). Here, the boundaries with irregular serrated configuration consist still of the single phase, as marked by the letter H, and the lamellar congeries with smaller size and various orientations exist in the interior of grain, as marked in the region J. The microstructure in the region C near fracture is shown in Fig. 7(d), which consists still of the regular lamellar/phases, the grain boundaries are located in between the lamellar / congeries with various orientations. Because the bigger plastic deformation occurs in the region, the original smooth and level lamellar / phases had been contorted to display the tortuous configuration, as marked by the arrow in Fig. 7(d), and the lamellar /phases in the region near boundaries display the disorder configuration, as marked in the region L. After crept up to fracture of the forged alloy at 900 oC/120MPa, the microstructure in the region near fracture are shown in Fig. 8. It is shown that the original lamellar phases in some regions still keep the smooth and level lamellar configuration, as marked by the letter K in Fig. 8(a), and the lamellar phases in different grains display the various orientations. If the lamellar congeries with various orientations are defined as different grains, the lamellar congeries in the grains display the various spacing distances. The regions between the lamellar congeries with different orientations are defined as grain boundaries which consist of fine block-like phases, as marked by the letter M in Fig. 8(a), and the size of boundaries in width increase to about 10-20m. Moreover, some needle-like precipitates are distributed in the lamellar phases and boundary regions, as marked by white arrow in Fig. 8(a). Fig. 8 Due to the bigger plastic deformation occurring in another region, the lamellar phases in some grains displays the distorted configuration, as shown in the regions N, P and Q of Fig. 8(b). Moreover, the fine block-like phases are distributed in the boundary regions, as marked by the thick arrow, and the crack appears in the boundary region which consists of fine block-like phases, as marked by the fine arrow, which may be thought to be the initiation of crack along boundaries. As the creep goes on, the crack is propagated along boundaries to form the bigger crack, as marked by the longer arrow in Fig. 8(c), and a smaller crack in the boundary region is marked by the fine arrow. When the bigger crack is further propagated along the boundary as the creep continues, the one may connect with the smaller crack to promote the propagation of crack up to the occurrence of creep fracture, which is thought to be the fracture mechanism of the forged alloy during creep. Therefore, it may be concluded that the boundaries in the forged alloy are still the weaker region of creep resistance during creep at high temperature. After crept for 237 h up to rupture under the applied stress of 120 MPa at 900 ºC, the microstructure of as-cast TiAl-Nb alloy are shown in Fig. 9, significant amount of dislocation rows appear in the local region of lamellar phases, the dislocation lines in the dislocation rows are parallel to each other, as marked by the black arrow in Fig. 9(a). The dislocation loop appears in the dislocation rows, as marked by the white arrow, the magnified morphology of the loop is indicated in the right-top of Fig. 9(a). Moreover, the regular dislocation rows distribute in the interface of the lamellar /phases, as shown in the square frame of Fig. 5
9(a), the magnified morphology of the dislocation rows is shown in the left down of Fig. 9(a). Moreover, some white granular particles are precipitated in the matrix, as marked by the short arrow in Fig. 9(a). In local region of sample, the lamellar phases are parallel to each other, as shown in Fig. 9(b), the size of the lamellar/phases is about 500—700m. Some dislocations have sheared into the lamellar phases, as marked by the black arrow in Fig. 9(b). Moreover, significant amount of dislocations are distributed in the interfaces of the lamellar phases to form the regular dislocations networks, as shown in the square frame, the magnified morphology in the square frame region is shown in the right down of Fig. 9(b). It is thought by analysis that the creep dislocations moving to the interface of phases may react with the dislocation network to change the original moving direction, which may promote the climbing of dislocations. Therefore, the dislocations networks located in the interfaces have the coordinating role on the deformation hardening and recovery softening of alloy during creep. Fig. 9 In another local region of sample, dislocation shearing into lamellar phases is marked by the arrow in Fig. 9(c), the dislocations located in the middle region of the photo display the dislocation rows configuration with parallel feature. Furthermore, the inclined boundary is located in between the lamellar phases, and the isoclinic stripes appearing in the inclined boundary are parallel to the orientation of lamellar phases. Some dislocations shear through the interface of lamellar phases, as shown in the square frame, and the magnified morphology in the square frame region is shown in the right top of Fig. 9(c). The fact that the creep dislocations within lamellar phases stop in the interface of phases indicates that the phase boundaries may hinder the slipping of dislocations. Therefore, it may be concluded that the deformation mechanism of the alloy during creep is dislocations slipping in the phases, wherein, the dislocations slipping in the matrix may react each other to form the dislocations networks, which may promote the climbing of dislocations. 3.4 Initiation and propagation of crack After as-cast TiAl-Nb alloy is crept for 237 h up to rupture at 900 ºC/120 MPa, the morphologies of the cracks initiating and propagating along boundaries near fracture are shown in Fig. 10, in which the direction of the applied stress is marked by the arrows. The boundary in the region is about 45 angles relative to the stress axis, as marked by straight line in Fig. 10(a), the orientation of the lamellar phases in the down side of the boundary is parallel to the direction of the boundary, but the orientation of the lamellar phases in the upper region of the boundary is perpendicular to the stress axis. Wherein, the cavity formed along the boundary of alloy during creep is marked by the fine arrow, which is regarded as the initiation of crack, the bigger crack along the boundary is regarded as the propagation of the crack, as marked by the bigger arrow in Fig. 10(a). It is understood from Fig. 10(a) that the propagating direction of the crack on one side is parallel to the orientation of the lamellar phases. In another region near fracture, the bigger crack along the boundary is still parallel to the orientation of the lamellar phases, and the boundary is at about 45 angles relative to the stress axis, as marked by the arrow in Fig. 10(b), but the propagating direction of the crack in another side is perpendicular to the orientation of the lamellar phases. This indicates that, during creep, the crack in the alloy is easily initiated and propagated along the boundaries at about 45 angles relative to the stress axis, which is attributed to the 45 angles boundaries bearing a bigger shear stress and its weaker strength. Fig. 10 6
In another region near fracture, the fracture surface originating from the crack propagation displays the various configurations. On the upper surface of the crack, the propagating direction of crack is inclined with the orientation of the lamellar phases, but on the down surface of the crack, the propagating direction of crack is parallel to the orientation of the lamellar phases, asshown in Fig. 10(c). The down surface of the crack formed during propagation displays the smooth feature, which suggests that the propagation of crack experiences a smaller resistance. But the fracture surface of displaying non smooth and tearing edges appears in the upper surface of the crack, as marked by the thicker arrow in Fig. 10(c), which suggests that the propagation of crack experiences a bigger resistance. It is thought by analysis that the microstructure of TiAl–Nb alloy consist of lamellar 2 phases, compared with Ti3Al, the strength of phase is weaker, therefore, the Ti3Al is the strengthening phase of the alloy. The grain boundary consisting of single phase is the weaker region of strength during creep of alloy, moreover, the boundaries at about 45 angles relative to the stress axis bear the bigger shear stress under loading. Therefore, the initiation and propagation of crack occurs easily in the boundaries along the direction at about 45 angles relative to the stress axis during creep. Furthermore, the resistance of the crack propagation is related to the orientation of the lamellar phases. When the direction of the crack propagation is parallel to the orientation of the lamellar phases, the propagation of crack along the single phase with weaker strength encounters only a smaller resistance. Therefore, the fracture of alloy after crept up to rupture displays a smooth surface, as marked by the fine arrow in Fig. 10(c). When the direction of the crack propagation is inclined with the orientation of the lamellar phases, the propagation of crack may encounter the Ti3Al phase with excellent strength to increase the resistance of crack propagation during creep, which results in the tearing edges appearing in the surface of fracture, as marked by the thicker arrow in Fig. 10(c). This is well consistent with the experimental results above state. 4. Discussion 4.1 Theory analysis of creep resistance of alloy at various states The microstructure of as-cast TiAl-Nb alloy consists of lamellar phases with different orientations, as shown in Fig. 1, wherein, the grain boundaries of alloy consists of single phase with smaller size in thickness. Compared to phase, the phase possesses a better strength, therefore, the deformation of the alloy during creep occurs mainly in phase. And it is thought that the characteristic of phases being easily deformed may delay the stress concentration resulting from the inharmonious deformation of the lamellar phases during creep. During isothermal forging, the bigger plastic deformation occurs in the alloy, which diminishes the grain size of them from 500m to about 80-120m. Therefore, compared to as cast TiAl-Nb alloy, the alloy after isothermal forged increases the quantity and volume fraction of the grain boundaries to a great extent. Especially, the wider boundary of the forged alloy consists of fine block-like phases with irregular features. During creep, the resistance of dislocations slipping in the boundary regions is smaller, the stress concentration originating from the large number dislocations piled up in the regions near boundary may results in the initiation and propagation of cracks along the boundaries, as shown in Fig. 8(b) and (c). Therefore, the wider boundaries of alloy are still the weaker regions of strength during creep at high temperature. It is thought by analysis that on the one hand, the bigger plastic deformation during isothermal forged refines the grains size of alloy, on the other hand, the process makes the original smooth and level lamellar phases transforming into the distorted configuration. Namely, the deformation of alloy during forging results in both decreasing creep resistance of microstructure and increasing the volume fraction of boundaries with weaker strength, so that the probability of the crack initiating and propagating along the
7
boundaries during creep at high temperature increases, which is thought to be the main reason of the forged alloy displaying a lower creep resistance and shorter lifetime. 4.2 Microstructure evolution and deformation features during creep The microstructure of as-cast TiAl-Nb alloy consists of lamellar /2 phases with different orientations, the size of lamellar phases in thickness is about 50-100 nm, as shown in Fig. 2, and the lamellar phases keep the coherent interface due to no interfacial dislocations distributed in between them. But crept for 237h up to rupture under the applied stress of 120 MPa at 900 ºC, the size of the lamellar phases in thickness increases to about 500-800 nm, as shown in Fig. 9(b), which suggests that the size of the lamellar phases in thickness increases as the creep goes on. Although the lamellar phases before creep keep the coherent interface, the interfacial dislocations have appeared in between the lamellar phases after crept for 237 h at high temperature, as marked by the square frame in Fig. 9(a) and (b). This indicates that the coherent interface in the lamellar phases has transformed into the semi-coherent interface as the size of the lamellar phases increases. It is thought by analysis that compared to the coherent interface, the extent of the lattice strain in the interface of phases decreases after the one is transformed into the semi-coherent interface, therefore, the resistance of dislocations in one phase shearing into another phase diminishes. Compared with Ti3Al phase, the phase possesses a weaker strength [22], so the plastic deformation of alloy during creep occurs firstly in the phase with weaker strength. Significant amount of dislocations may firstly activate in phase, a few dislocations appears in phase, as shown in Fig. 9(a) and (b). And the quantity of dislocations shearing into the lamellar phase increases with the strain of alloy, wherein, the dislocations both shear into the phase in the form of the rows and slip in the phase in the form of ring, as marked by the arrow in Fig. 9(a) and (b). As the creep goes on, significant amount of dislocations slipping to the interface are hindered to pile up in the region near the interface, which may bring the stress concentration. The dislocations piled up in phase may shear into the phase along the interface dislocation network when the value of the stress concentration exceeds the strength of 2 phase. Moreover, significant amount of dislocations shearing into the lamellar phases are alternately activated, which may twist the lamellar phases within grains, as marked by the white arrow in Fig. 7(c), to increase the strain of alloy. Moreover, the grain boundaries in between the lamellar congeries display the irregular serrated configuration, and the ones consist of the single phase with Ll0 tetragonal structure [21], as marked by the arrow in Fig. 1(a). Because the TiAl phase possesses a weaker strength and better ductility, the bigger plastic deformation occurs in the boundary region of as-cast alloy during creep at high temperature. The boundary regions with wider feature are marked in the regions H and J of Fig 7(c), which suggests that the grain boundary is a weaker link of restricting the creep strength of alloy. It is indicated from Fig. 8 and 10 that the initiation and propagation of crack occurs mainly in the grain boundaries at about 45 angles relative to the stress axis during creep of alloy due to the direction bearing a bigger shear stress. It is thought by analysis that the boundaries of as cast and forged alloys consist of single phase and fine block-like phases with weaker strength, during creep at high temperature, the bigger shear stress may promote the initiation of crack along the boundaries with weaker strength. And as the creep goes on, the bigger shear stress promotes the propagation of crack along the boundaries up to the occurrence of creep fracture. Fig. 11 The schematic diagram of crack initiated and propagated along the boundary at 45 angles relative to the stress axis is shown in Fig. 11, the direction of the applied stress is marked by the arrows. Because the boundaries of alloys possess a weaker strength, the initiation of the crack occurs firstly in the inclined 8
boundary under the action of the bigger shearing stress, which is schematically shown in Fig. 11(a). As the creep goes on, the size of the crack propagating along boundary increases under the action of bigger shearing stress, as shown in Fig. 11(b). And then the crack is further propagated along the inclined boundary until encountering the boundary with another orientation, as marked by the arrow in Fig. 10(b), which may hinder the propagation of the crack, as schematically shown in Fig. 11(c). Although the grain boundary is the weaker link of strength during creep of the as cast alloy, the propagation of the cracks along boundary may encounter the Ti3Al phase with better strength to hinder the one of them. Therefore, when the orientation of the lamellar phases is inclined with the direction of the grain boundaries, the inclined lamellar Ti3Al phase relative to the boundary may restrain the propagation of the cracks to improve the creep resistance and life of as cast alloy. The analysis stated above is well agreement with the experiment results. 4. Conclusion (1) The microstructure of as-cast TiAl-Nb alloys consists mainly of lamellar 2 phases, the boundaries with irregular serrated configuration consist of single phase, and the ones are located in between the lamellar 2 phases with various orientations. (2) The bigger plastic deformation of alloy during isothermal forging may refine the size of the grains and increase the volume fraction of boundaries which consists of the fine block-like 2 phases with weaker strength. During creep of the forged alloy, the wider size and bigger volume fraction of boundaries increase the probability of crack initiated and propagated along boundaries, which is thought to be the main reason of the forged alloy displaying a lower creep resistance and shorter lifetime. (3) During creep at temperature near 900 oC, compared to forged alloy, the as-cast alloy displays a better creep resistance and longer creep life. The deformation mechanism of alloys during creep is significant amount of dislocations shearing into the lamellar phases in the form of dislocations rows and rings. Moreover, the dislocations networks located in the interfaces may promote the climbing of dislocations to retard the stress concentration and improve the creep resistance of alloy. (4) Compared with Ti3Al phase, the phase possesses a weaker strength, therefore, the crack is easily initiated and propagated along the boundaries at about 45 angles relative to the stress axis until the occurrence of creep fracture. Wherein, the fracture parallel to the lamellar phases displays the smooth surface, while the tearing edges appear on the surface of fracture in one side of the inclined boundaries relative to lamellar phases, which is attributed to the Ti3Al phase with better strength hindering the propagation of crack during creep. Acknowledgements Sponsorship of this research by the National Basic Research Program of China under Grant No. 2011CB605506 is gratefully acknowledged. References [1] Zheng Z, Jiang Y, Dong HX, Tang LM, He YH, Huang BY. Environmental corrosion resistance of intermetallic compounds. Trans. Nonferrous Met Soc China 2009; 19: 582-585. [2] Ye XC, Su YQ, Guo JJ, Zhang LH, Hao ZJ. Effect of heat treatment on microstructure of TiAl based alloy casting by suction casting process, Trans Mater Heat Treat 2012; 33: 132-135. [3] Lu X, Zhao LH, Zhu LP, Zhang B, Qu XH. High-temperature mechanical properties and deformation behavior of high Nb containing TiAl alloys fabricated by spark plasma sintering. Mater 2012; 19(4): 354-359. 9
[4] Ding XF, Lin JP, Zhang LQ, Su YQ, Chen GL. Microstructural control of TiAl–Nb alloys by directional solidification. Acta Mater 2012; 60(2): 498-506. [5] Hénaff G, Gloanec AL. Fatigue properties of TiAl alloys. Intermetallics 2005; 13: 543-558. [6] Lin JP, Xu XJ, Wang YL, He SF, Zhang Y, Song XP, Chen GL. High temperature deformation behaviors of a high Nb containing TiAl alloy. Intermetallics 2007; 15: 668–674. [7] Kong FT, Chen YY, Tian J, Chen ZY. Development in TiAl Based intermetallics research. Master Sci Technol 2003; 11: 441-444. [8] Chen GQ, Zhang BG, Liu W, Feng JC. Crack formation and control upon the electron beam welding of TiAl-based alloys. Intermetallics 2011; 19: 1857-1863. [9] Kelly TJ, Juhas MC, Huang SC. Effect of a B2/gamma structure on the tensile properties of the cast gamma titanium aluminide Ti-48Al-2Cr-2Nb. Scr Metall Mater 1993; 29: 1409-1414. [10] McCullough C, Valencia JJ, Levi CG, Mehrabian R. Phase equilibria and solidification in Ti-Al alloys. Acta Metall Master 1989; 37: 1321-1336. [11] Naka S, Thomas M, Sanchez C, Khan, T. Development of third generation role of solidification paths. TMS 1997; 313-322. [12] Sun FS, Cao CX, Yan MQ, Kim SE, Yang TL. Alloying mechanism of beta stabilizers in a-TiAl alloy. Metall Mater Trans A 2001; 32: 1573-1589. [13] Liu ZC, Lin PJ, Chen GL. Effect of the addition W on the microstructure and mechanical properties for high-Nb TiAl alloy. Trans Mater Heat Treat 2001; 22: 8-13. [14] Hu D, Yang C, Huang A, Dixon M, Hecht U. Grain refinement in beta-solidifying Ti44A18Nb1B. Intermetallics 2012; 23: 49-56. [15] Lin JG, Zhang YG, Chen CQ. Effects of the microstructure of γ-TiAl alloys on the creep behavior. Journal of Beijing University of Aeronautics and Astronautics 1998; 24: 734-737. [16] Chen YY, Kong FT, Han JC, Chen ZY, Tian J. Influence of yttrium on microstructure, mechanical properties and deformability of Ti-43A1-9V alloy. Intermetallics 2005; 13: 263-266. [17] Yan YQ, Zhou L, Wang WS, Zhang YN. 8.5 Nb–TiAl alloy with fine grains. Journal of alloys and compounds 2003; 361: 241-246. [18] Wang YH, Lin JP, He YH, Wang YL, Chen GL. Diffusion behavior of Nb element in high NB containing TiAl alloys by reactive hot pressing. Rare Met 2006; 25: 349-354. [19] Kim HY, Maruyama K. Changes in lamellar microstructure by parallel twinning during creep in soft PST crystal of TiAl alloy. Trans Nonferrous Met Soc China 2002; 12: 561-568. [20] Niu HZ, Chen YY, Xiao SL, Xu J. Microstructure and mechanical poperties of a novel beta g-TiAl alloy. Intermetallics 2012; 31: 225-231. [21] Qu HP, Li P, Zhang SQ, Li A, Wang HM. The effects of heat treatment on the microstructure and mechanical property of laser melting deposition γ-TiAl intermetallic alloys. Materials and Design 2010; 31: 2201-2210. [22] Li XW, Sun HF, Zhang P, Fang WB. The effect of strain on dynamic recrystallization of PM Ti–45Al10Nb intermetallics during isothermal forging. Intermetallics 2014; 55: 90-94. [23] Tian SG, Wang Q, Yu HC, Sun HF, Li QY. Microstructure and creep behaviors of a high Nb-TiAl intermetallic compound based alloy. Mater. Sci Eng. A 614, 2014: 338-346.
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(a)
(b)
40m
20m
Fig. 1 Microstructure of the as-cast TiAl-Nb alloy. (a) macro-morphology, (b) magnified morphology.
Fig. 2 After heat treatment, microstructure of TiAl-Nb alloy.
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Fig. 3 Microstructure of the isothermal forged TiAl-Nb alloy.
Strain,ε(%)
20 15
1 -- As-cast alloy 2 -- Forged alloy 120 MPa o T -- 900 C
10 5
0
2
1
50
100 150 200 250 Time, (h) Fig. 4 Creep curves of alloy with various states under the applied stress of 120MPa at 900 ℃
12
30
30
18
24 Strain,
Strain,ε(%)
24
-- 140 MPa
1 -- 120 MPa 2 -- 130 MPa 3 -- 140 MPa o T -- 900 C
12 3
2
6
18
1 -- 890 oC o 2 -- 900 C o 3 -- 910 C
12 3
1
2
6
1
(b)
(a) 0
50
100 150 Time, (h)
200
0
250
20
40
60 80 Time, (h)
100
120
Fig. 5 Creep curves of alloy at different conditions. (a) Applied different stresses at 900 ºC, (b) applied stress of 140MPa at different temperatures.
-3.30
-3.3
Q = 413.9 kJ/mol
-3.45
-3.6
ln ss
ln ss
-3.60 -3.75 -3.90 -4.05 -4.20
n = 7.80
-3.9 -4.2 -4.5
(a) 8.46
(b) 8.49
8.52 8.55 1/T,10-3K-1
8.58
8.61
-4.8
4.80
4.84
4.88
4.92
4.96
In
Fig. 6 Dependence of the strain rates of alloy during steady state creep on temperatures and stresses. (a) Strain rates and temperatures, (b) strain rates and applied stress
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Fig. 7 Microstructure in different regions of sample after crept up to fracture at 120 MPa/900 ºC. (a) Schematic diagram of marking observed regions in sample, (b), (c), (d) being SEM morphologies corresponding to A, B, C regions, respectively.
Fig. 8 Microstructure in different regions of the forged sample after crept up to fracture at 120 MPa/900 ºC. (a) lamellar congeries in different grains, (b) crack initiated along boundaries, (c) the propagation of crack.
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Fig. 9 Microstructure of alloy crept for 237h up to fracture at 900 ºC/120 MPa. (a) dislocation rows, (b) interfacial dislocations networks, (c) dislocations shearing lamellar phase.
Fig. 10 After crept for 237 h up to fracture at 900oC/120MP, initiation and propagation of cracks. (a) Initiation of crack along the boundary at about 45 angles relative to the stress axis, (b) propagation of crack along the boundary, (c) non smooth surface formed in another side of crack
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Fig. 11 Schematic diagram of crack initiated and propagated along boundary at 45 angle relative to the stress axis during creep at 900 oC. (a) Initiation of crack along boundary, (b) propagation of crack along boundary, (c) propagation of crack being stopped in the boundary with another orientation,
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