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Microstructure degradation in high temperature fatigue of TiAl alloy
International Journal of Fatigue xxx (2013) xxx–xxx
Contents lists available at ScienceDirect
International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Microstructure degradation in high temperature fatigue of TiAl alloy T. Kruml a,⇑, K. Obrtlík b a b
CEITEC IPM, Zizkova 22, Brno, Czech Republic Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, Brno, Czech Republic
a r t i c l e
i n f o
Article history: Received 22 March 2013 Received in revised form 26 August 2013 Accepted 9 October 2013 Available online xxxx Keywords: Low cycle fatigue Lamellar TiAl alloy High temperature fatigue Dislocations Microstructure
a b s t r a c t Low cycle fatigue properties of lamellar TiAl with 8 at.% Nb were studied at four temperatures: room temperature, 700, 750 and 800 °C. Up to 750 °C, stable cyclic behaviour is observed while cyclic softening is characteristic for 800 °C. The strength of the alloy is still high even at 800 °C. The TEM observation did not reveal any substantial changes in the microstructure due to the cycling at RT. At 750 °C, the lamellar structure was in some places destroyed by cyclic plastic straining and pure c-phase islands with high density of dislocation debris were formed. At 800 °C, the domains without lamellar structure cover about 10% of volume and are almost dislocation free. The destruction of lamellar microstructure and possible annealing of dislocation debris is the reason for marked cyclic softening at 800 °C. Ó 2013 Published by Elsevier Ltd.
1. Introduction The development of materials derived from binary TiAl alloys continues for at least 35 years. Thanks to the high temperature strength, good corrosion resistance and low density, the main domain of possible applications concerns parts subjected to loads at elevated temperatures. Indeed, lamellar TiAl alloys recently have started being used for car turbochargers or parts of plane engines. Fatigue properties of such components working in rotation and subjected to vibrations and temperature gradients are evidently key characteristics necessary for engineering design. Since the pioneering work of Sastry and Lipsitt [1], fatigue properties of this class of materials were regularly studied. Most of these effort was devoted to binary Ti–Al alloys [2,3] and later to alloys with 2 at.% of Nb, e.g. by Hénaff et al. [4–6]. The interest of the alloy with 2Nb is understandable since this variant of TiAl alloys is already used in several applications [5,7]. It can be summarised that, in general, TiAl alloys show behaviour typical for brittle materials: large scatter of fatigue life, strong influence of material defects, strong influence of surface finishing [8], flat S–N curves resulting in fatigue limit close to the uniaxial yield point [4], usually no surface slip band formation (exceptions are reported in [3,9,10]) and fast crack propagation with Paris exponents 5–10 times higher than typical values of metallic systems [11,12]. Because 2 at.% of Nb addition improves the performance of the material considerably, alloys with even higher amount of Nb were
⇑ Corresponding author. Tel.: +420 532 290 379. E-mail address: [email protected] (T. Kruml).
tested. Alloys with 5–10 at.% of Nb are called the third generation of TiAl alloys. Much less data about low cycle fatigue (LCF) behaviour of these alloys are reported up to now [11,13]. It has been documented that increase of Nb amount to 8 at.% further improves the strength of material: the yield stress is higher by 35% for the 8Nb alloy at RT and by 25% at 750 °C. The fracture stress at room temperature is also higher by about 180 MPa for the 8Nb alloy due to both higher strength and higher fracture plastic strain [14]. Substantial increase of fatigue life in comparison with 2Nb alloy in S–N diagram alloy was reported [15]. Roth and Biermann [16] compared data obtained on 5Nb alloy fatigued up to 800 °C in constant strain amplitude tests LCF and thermomechanical fatigue (TMF) tests. They found that the fatigue life in both LCF and TMF regimes can be well predicted by Smith, Watson and Topper damage parameter [17]. In this paper, results of systematic study of cyclic properties and fatigue damage mechanisms in TiAl–8 at.%Nb are described at room temperature (RT), 700 °C, 750 °C and 800 °C.
2. Experimental procedure 2.1. Material The material was supplied in the form of a cylindrical ingot. It was prepared by casting in the Flowserve company. The chemical composition of the alloy is shown in Table 1. Metallographic inspection revealed differences in the microstructure between the centre and borders of the ingot. Therefore, the cylindrical fatigue specimens were manufactured only from the outer regions of
0142-1123/$ - see front matter Ó 2013 Published by Elsevier Ltd. http://dx.doi.org/10.1016/j.ijfatigue.2013.10.005
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
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T. Kruml, K. Obrtlík / International Journal of Fatigue xxx (2013) xxx–xxx
Table 1 Chemical composition of the TiAl alloy. Element
Ti
Al
Nb
Cr
Ni
Si
B
at.%
47.8
44.2
7.8
0.7
0.2
0.1
–
the ingot. It was observed that the scatter of the fatigue data was relatively low which is probably linked to the careful specimen selection and surface preparation. The microstructure of the material is lamellar with some single c-phase islands close to grain boundaries covering of about 3% volume. Minority phases b and B2 were observed either in regions between lamellas of c and a2 phases or in regions close to grain boundaries. Fine lamellar structure of c and a2 phases is shown in Fig. 1. The c phase lamellas have average thickness of 0.87 lm while a2 lamellas are thinner, typically less than 100 nm in thickness. The monotonic tensile 0.1% proof stress of the material at RT and 750 °C is 548 and 449 MPa, respectively. 2.2. Low cycle fatigue tests Cylindrical specimens with 6 mm in diameter and 15 mm of gauge length were prepared. The material was ductile enough to make standard machining operation as turning, grinding and drilling. The gauge length was mechanically and then electrolytically polished. Fatigue tests were performed under strain control using MTS servohydraulic machines in symmetric tension–compression cycle (Re = 1). The total strain amplitude and strain rate of
2 10 3 s 1 were kept constant in all tests. Strain was measured by extensometers attached directly to the gauge part of the specimens. The tests were performed at RT, 700 °C, 750 °C and 800 °C in air. Heating was provided by a three-zone resistant furnace and monitored by three thermocouples attached to both specimen ends and to the upper part of the gauge section. 2.3. Microstructural observations The microstructure was observed on the polished and etched surface by light microscopy and scanning electron microscopy (Jeol 6460). Thin foils for transmission electron microscopy were prepared parallel, perpendicular and with the angle of 45° to the loading axis. Philips CM12 operating at 120 kV was used.
3. Results 3.1. Cyclic behaviour Cyclic hardening/softening curves are shown in Fig. 2a–d. From RT to 750 °C, almost stable cyclic behaviour during testing with strain amplitudes from 0.25% to 0.4% is observed. Noticeable cyclic softening appears at 800 °C, especially at high strain amplitudes. Since cycling does not alter the properties of the material at RT, tensile curve and cyclic stress–strain curve coincide (Fig. 3). Excellent behaviour of the material in cycling at high temperatures is documented in Fig. 3. The stress amplitude at half-life at given strain amplitude decreases slowly with temperature. Even at 800 °C, stress amplitudes are still above 300 MPa. While the level of stress amplitude for the three high temperatures is similar in Fig. 3, clear difference of derived Wöhler curves (stress amplitude at half-life versus number of cycles to fracture NF) appears for every testing temperature (Fig. 4). Experimental data were fitted with a polynomial regression to show the course of the fatigue life curves at different temperatures relative to each other. The main detrimental effect of the increase of testing temperature from 700 to 800 °C is thus the shortening of fatigue life. It is not known if the shortening of the fatigue life is due to the faster fatigue crack growth or faster crack initiation, or both. It should be noted that S–N curves measured in strain controlled test and stress control test coincide only for cyclically stable material which is not fulfilled at 800 °C. Fatigue life in relation to strain amplitude ea, which was kept constant during cycling, is shown in Fig. 5. In low cycle fatigue domain (NF < 104) where the plastic part of the strain is high, the material shows shortest life at RT due to its room temperature brittleness. For lower ea, the fatigue life is shortest for specimens cycled at 800 °C. The difference between the fatigue life at 750 and 800 °C is significant especially for high cycle fatigue regime. Plotting the fatigue life data against plastic strain amplitude eap at half-life (Fig. 6) reveals a surprising effect. While at RT the usual Coffin–Manson power law is, taking into account the data scatter, valid for all plastic strain amplitudes with the exponent c = 0.25, two different regimes are observed for high temperature data. In the low cycle fatigue domain, the data can be fitted by the power law with c around 0.3. However, the curve at Fig. 6 is almost horizontal for NF > 104 for all three high testing temperatures. 3.2. Microstructure evolution in cyclic deformation
Fig. 1. Lamellar microstructure of the material. (a) TEM bright field and (b) TEM HAADF mode, white lamellas are the a2 phase.
The changes in the microstructure after fatigue failure were studied in several specimens cycled at various temperatures and strain amplitudes. It can be concluded that at RT, no changes in
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
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Fig. 2. Cyclic hardening/softening curves. Testing temperature and strain amplitude are given in inserts.
Fig. 3. Tensile curve at RT and cyclic stress–strain curves at the four testing temperatures.
the internal structures were detected. This observation agrees with the stable cyclic response of the material. In specimens subjected to fatigue to failure at 750 °C, a microstructure appears which has never been observed in as-received state. It consists of areas of c-phase inside the grains, without lamellas of a2 phase or c-phase variants, with high dislocation density and a lot of dislocation debris in the form of small prismatic loops (Fig. 7). Formation of such structure is a consequence of cyclic plastic loading at 750 °C. The described areas are quite rare, the relative volume of destroyed lamellar structure is therefore very small. For this reason, this microstructural change is not visible on the cyclic hardening/softening curves at Fig. 2b.
Fig. 4. S–N curves at the four testing temperatures.
At 800 °C, the destruction of lamellar structure due to cyclic loading is much more frequent than at 750 °C. Contrary to the observation of specimens cycled at 750 °C, prismatic loops are not observed and dislocation density in the newly formed c phase is low (Fig. 8), substantially lower than at 750 °C (compare Figs. 7b and 8b). These c-phase islands cover approximately 10% of volume of the specimens.
4. Discussion TiAl alloys at RT usually cyclically harden [4,18]. In particular, very strong cyclic hardening was reported for alloys with 2 at.%
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
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Fig. 5. Combined fatigue life curve with guidelines.
Fig. 7. (a) Partially destroyed lamellar microstructure; and (b) high dislocation density with prismatic loops in the c phase. 750 °C, ea = 0.41%, ra = 443 MPa, foil plane parallel to the loading axis, NF = 438 cycles.
Fig. 6. Coffin–Manson diagram with guidelines.
Nb or 2Nb–2Cr [4,19]. It was explained by intensive twinning. In the present material, no cyclic hardening was measured and twins were observed only rarely. It is thus possible that the explanation of twinning induced cyclic hardening is correct and that higher amount of Nb reduces the twinning activity, e.g. by alternating the stacking fault energy. Concerning the differences of the cyclic behaviour at different testing temperatures, the changes in material response appear between 750 and 800 °C. Increase of testing temperature to 800 °C shortens the fatigue life considerably in S–N representation and material cyclically softens. This means that the maximum working temperature for applications demanding high strength of material is around to 750 °C. The change of the usual slope of Coffin–Manson diagram in low cycle fatigue domain to almost horizontal line for NF > 104 has, up to our knowledge, never been reported. On the contrary, inverse change of the slope in the Coffin–Manson diagram was observed for lamellar Ti–46.6Al–1.4Mn–2Mo cycled in air at 800 °C [20]: c = 0.22 for NF < 7 104 and c = 0.65 for NF > 7 104. Authors explained this change of the slope by change of cracking mode from transgranular to intergranular due to a2 ? c transformation on grain boundaries. In our case, intergranular cracking was not observed. However, two slopes in the Coffin–Manson diagram,
detailed for 750 °C in Fig. 9, suggests two operating mechanisms of cyclic straining whose participation depends on strain amplitude. Additional data, particularly relative to the microstructure of TiAl–8Nb alloy are necessary to explain the behaviour. The main result of the TEM microstructural observation is the documentation of changes of the microstructure of specimens cycled at 800 °C: in substantial part of the volume, lamellar microstructure is replaced by islands of c phase with low dislocation density. The destruction of lamellar microstructure occurs by gradual disappearance of a2 lamellas. As seen in Fig. 8a, in areas of 10 lm the a2 lamellas disappeared, leaving bands of dislocations on their initial places (Fig. 8b). It is supposed that with further cycling, the process would continue and resting a2 lamellas would continue to shorten, due to combined effect of temperature and cyclic straining. The disappearance of lamellar microstructure during cyclic deformation at high temperatures has already been observed by Appel et al. [11]. The microstructure of the specimen after fatigue at 850 °C published in [11] closely resembles Fig. 7a. In this work, the onset of destruction of the lamellar microstructure due to cyclic loading was assigned to 750 °C. An interesting difference between c islands formed at 750 °C and 800 °C is the different level of dislocation density including prismatic loops, which are abundant at 750 °C and non-existing at 800 °C. It can be speculated that diffusion processes, necessary for the annealing of prismatic loops, are much faster at 800 °C than at 750 °C.
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Stable cyclic behaviour was observed up to 750 °C, clear cyclic softening appears at 800 °C. Fatigue strength of the material is still high at 800 °C (cyclic yield stress above 300 MPa). With increasing temperature, significant reduction of fatigue life is observed in S–N diagram, particularly between 750 and 800 °C. Destruction of lamellar microstructure was observed at 750 °C and, more pronounced, at 800 °C. Areas of single c phase are formed. High density of dislocations and prismatic loops are observed at 750 °C while low dislocation density is found in these areas at 800 °C. Accentuated destruction of lamellas due to cyclic loading is responsible for cyclic softening at 800 °C.
Acknowledgements Support of the project of the Czech Science Foundation 107/11/ 0704 and by the research Project No. AV0Z 20410507 of the Academy of Sciences of the Czech Republic is acknowledged. The research was conducted in CEITEC research infrastructure supported by the project. References
Fig. 8. (a) Partially destroyed lamellar microstructure; and (b) the front of the resting lamellar microstructure and remaining dislocations in the c phase. 800 °C, ea = 0.24%, ra = 306 MPa, foil plane parallel to the loading axis, NF = 61,174 cycles.
Fig. 9. Coffin–Manson diagram for 750 °C. Dashed line is the extrapolation of data from the high cycle to low cycle domain.
5. Conclusions Low cycle fatigue behaviour of TiAl alloyed with 7.8 at.% Nb was studied from RT to 800 °C.
[1] Sastry SML, Lipsitt HA. Fatigue deformation of TiAl base alloys. Metall Trans A 1977;8A:299–308. [2] Umakoshi Y, Yasuda HY, Nakano T. Plastic anisotropy and fatigue of TiAl PST crystals: a review. Intermetallics 1996;4:65–75. [3] Yasuda H, Nakano T, Umakoshi Y. Cyclic deformation-behavior of Ti–Al alloys containing oriented lamellae. Philos Mag A 1995;71:127–38. [4] Hénaff G, Gloanec A-L. Fatigue properties of TiAl alloys. Intermetallics 2005;13:543–58. [5] Gloanec A-L, Jouiad M, Bertheau D, Grange M, Hénaff G. Low-cycle fatigue and deformation substructures in an engineering TiAl alloy. Intermetallics 2007;15:520–31. [6] Jouiad M, Gloanec A-L, Grange M, Hénaff G. Cyclic deformation mechanisms in a cast gamma titanium aluminide alloy. Mat Sci Eng A 2005;400–401:409–12. [7] Badami M, Marino F. Fatigue tests of un-HIPed gamma-TiAl engine valves for motorcycles. Int J Fatigue 2006;28:722–32. [8] Vaidya WV, Schwalbe KH, Wagner R. Understanding the fatigue resistance of gamma titanium aluminide. In: Kim Y-W, Wagner R, Yamaguchi M, editors. Gamma titanium aluminides. Warrendale: TMS; 1995. p. 867–74. [9] Petrenec M, Polák J, Bucˇek P. Cyclic plasticity and strain localization in cast cTiAl based alloy. Proc Eng 2011;10:1390–5. [10] Huang Z, Bowen P. Persistent microslip bands in the lamellar TiAl structure subjected to room temperature fatigue. Scr Mater 2001;45:931–7. [11] Appel F, Heckel TK, Christ H-J. Electron microscope characterization of low cycle fatigue in a high-strength multiphase titanium aluminide alloy. Int J Fatigue 2010;32:792–8. [12] Chan KS, Shih DS. Fundamental aspects of fatigue and fracture in a TiAl sheet alloy. Metall Mater Trans 1998;29A:73–87. [13] Cui WF, Liu CM. Fracture characteristics of c-TiAl alloy with high Nb content under cyclic loading. J Alloys Compd 2009;477:596–601. [14] Kruml T, Petrenec M, Obrtlík K, Polák J, Bucek P. Influence of niobium alloying on the low cycle fatigue of cast TiAl alloys at room and high temperatures. Proc Eng 2010;2:2297–305. [15] Kruml T, Dlouhy´ A, Petrenec M, Obrtlík K, Polák J. Fatigue properties of high Nb TiAl alloy. J Phys: Conf Ser 2010;240:012057. [16] Roth M, Biermann H. Thermo-mechanical fatigue behaviour of a modern cTiAl alloy. Int J Fatigue 2008;30:352–6. [17] Smith KN, Watson P, Topper TH. A stress–strain function for the fatigue of metals. J Mater 1970;5:767–78. [18] Satoh M, Horibe S, Nakamura M, Uchida H. Cyclic deformation and fatigue in TiAl intermetallic compound under plastic strain control. Int J Fatigue 2010;32:698–702. [19] Gloanec A-L, Milani T, Hénaff G. Impact of microstructure, temperature and strain ratio on energy based low-cycle fatigue life prediction models for TiAl alloys. Int J Fatigue 2010;32:1015–21. [20] Park YS, Nam SW, Hwang SK, Kim NJ. The effect of the applied strain range on fatigue cracking in lamellar TiAl alloy. J Alloys Compd 2002;335:216–23.
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
Contents lists available at ScienceDirect
International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Microstructure degradation in high temperature fatigue of TiAl alloy T. Kruml a,⇑, K. Obrtlík b a b
CEITEC IPM, Zizkova 22, Brno, Czech Republic Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, Brno, Czech Republic
a r t i c l e
i n f o
Article history: Received 22 March 2013 Received in revised form 26 August 2013 Accepted 9 October 2013 Available online xxxx Keywords: Low cycle fatigue Lamellar TiAl alloy High temperature fatigue Dislocations Microstructure
a b s t r a c t Low cycle fatigue properties of lamellar TiAl with 8 at.% Nb were studied at four temperatures: room temperature, 700, 750 and 800 °C. Up to 750 °C, stable cyclic behaviour is observed while cyclic softening is characteristic for 800 °C. The strength of the alloy is still high even at 800 °C. The TEM observation did not reveal any substantial changes in the microstructure due to the cycling at RT. At 750 °C, the lamellar structure was in some places destroyed by cyclic plastic straining and pure c-phase islands with high density of dislocation debris were formed. At 800 °C, the domains without lamellar structure cover about 10% of volume and are almost dislocation free. The destruction of lamellar microstructure and possible annealing of dislocation debris is the reason for marked cyclic softening at 800 °C. Ó 2013 Published by Elsevier Ltd.
1. Introduction The development of materials derived from binary TiAl alloys continues for at least 35 years. Thanks to the high temperature strength, good corrosion resistance and low density, the main domain of possible applications concerns parts subjected to loads at elevated temperatures. Indeed, lamellar TiAl alloys recently have started being used for car turbochargers or parts of plane engines. Fatigue properties of such components working in rotation and subjected to vibrations and temperature gradients are evidently key characteristics necessary for engineering design. Since the pioneering work of Sastry and Lipsitt [1], fatigue properties of this class of materials were regularly studied. Most of these effort was devoted to binary Ti–Al alloys [2,3] and later to alloys with 2 at.% of Nb, e.g. by Hénaff et al. [4–6]. The interest of the alloy with 2Nb is understandable since this variant of TiAl alloys is already used in several applications [5,7]. It can be summarised that, in general, TiAl alloys show behaviour typical for brittle materials: large scatter of fatigue life, strong influence of material defects, strong influence of surface finishing [8], flat S–N curves resulting in fatigue limit close to the uniaxial yield point [4], usually no surface slip band formation (exceptions are reported in [3,9,10]) and fast crack propagation with Paris exponents 5–10 times higher than typical values of metallic systems [11,12]. Because 2 at.% of Nb addition improves the performance of the material considerably, alloys with even higher amount of Nb were
⇑ Corresponding author. Tel.: +420 532 290 379. E-mail address: [email protected] (T. Kruml).
tested. Alloys with 5–10 at.% of Nb are called the third generation of TiAl alloys. Much less data about low cycle fatigue (LCF) behaviour of these alloys are reported up to now [11,13]. It has been documented that increase of Nb amount to 8 at.% further improves the strength of material: the yield stress is higher by 35% for the 8Nb alloy at RT and by 25% at 750 °C. The fracture stress at room temperature is also higher by about 180 MPa for the 8Nb alloy due to both higher strength and higher fracture plastic strain [14]. Substantial increase of fatigue life in comparison with 2Nb alloy in S–N diagram alloy was reported [15]. Roth and Biermann [16] compared data obtained on 5Nb alloy fatigued up to 800 °C in constant strain amplitude tests LCF and thermomechanical fatigue (TMF) tests. They found that the fatigue life in both LCF and TMF regimes can be well predicted by Smith, Watson and Topper damage parameter [17]. In this paper, results of systematic study of cyclic properties and fatigue damage mechanisms in TiAl–8 at.%Nb are described at room temperature (RT), 700 °C, 750 °C and 800 °C.
2. Experimental procedure 2.1. Material The material was supplied in the form of a cylindrical ingot. It was prepared by casting in the Flowserve company. The chemical composition of the alloy is shown in Table 1. Metallographic inspection revealed differences in the microstructure between the centre and borders of the ingot. Therefore, the cylindrical fatigue specimens were manufactured only from the outer regions of
0142-1123/$ - see front matter Ó 2013 Published by Elsevier Ltd. http://dx.doi.org/10.1016/j.ijfatigue.2013.10.005
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
2
T. Kruml, K. Obrtlík / International Journal of Fatigue xxx (2013) xxx–xxx
Table 1 Chemical composition of the TiAl alloy. Element
Ti
Al
Nb
Cr
Ni
Si
B
at.%
47.8
44.2
7.8
0.7
0.2
0.1
–
the ingot. It was observed that the scatter of the fatigue data was relatively low which is probably linked to the careful specimen selection and surface preparation. The microstructure of the material is lamellar with some single c-phase islands close to grain boundaries covering of about 3% volume. Minority phases b and B2 were observed either in regions between lamellas of c and a2 phases or in regions close to grain boundaries. Fine lamellar structure of c and a2 phases is shown in Fig. 1. The c phase lamellas have average thickness of 0.87 lm while a2 lamellas are thinner, typically less than 100 nm in thickness. The monotonic tensile 0.1% proof stress of the material at RT and 750 °C is 548 and 449 MPa, respectively. 2.2. Low cycle fatigue tests Cylindrical specimens with 6 mm in diameter and 15 mm of gauge length were prepared. The material was ductile enough to make standard machining operation as turning, grinding and drilling. The gauge length was mechanically and then electrolytically polished. Fatigue tests were performed under strain control using MTS servohydraulic machines in symmetric tension–compression cycle (Re = 1). The total strain amplitude and strain rate of
2 10 3 s 1 were kept constant in all tests. Strain was measured by extensometers attached directly to the gauge part of the specimens. The tests were performed at RT, 700 °C, 750 °C and 800 °C in air. Heating was provided by a three-zone resistant furnace and monitored by three thermocouples attached to both specimen ends and to the upper part of the gauge section. 2.3. Microstructural observations The microstructure was observed on the polished and etched surface by light microscopy and scanning electron microscopy (Jeol 6460). Thin foils for transmission electron microscopy were prepared parallel, perpendicular and with the angle of 45° to the loading axis. Philips CM12 operating at 120 kV was used.
3. Results 3.1. Cyclic behaviour Cyclic hardening/softening curves are shown in Fig. 2a–d. From RT to 750 °C, almost stable cyclic behaviour during testing with strain amplitudes from 0.25% to 0.4% is observed. Noticeable cyclic softening appears at 800 °C, especially at high strain amplitudes. Since cycling does not alter the properties of the material at RT, tensile curve and cyclic stress–strain curve coincide (Fig. 3). Excellent behaviour of the material in cycling at high temperatures is documented in Fig. 3. The stress amplitude at half-life at given strain amplitude decreases slowly with temperature. Even at 800 °C, stress amplitudes are still above 300 MPa. While the level of stress amplitude for the three high temperatures is similar in Fig. 3, clear difference of derived Wöhler curves (stress amplitude at half-life versus number of cycles to fracture NF) appears for every testing temperature (Fig. 4). Experimental data were fitted with a polynomial regression to show the course of the fatigue life curves at different temperatures relative to each other. The main detrimental effect of the increase of testing temperature from 700 to 800 °C is thus the shortening of fatigue life. It is not known if the shortening of the fatigue life is due to the faster fatigue crack growth or faster crack initiation, or both. It should be noted that S–N curves measured in strain controlled test and stress control test coincide only for cyclically stable material which is not fulfilled at 800 °C. Fatigue life in relation to strain amplitude ea, which was kept constant during cycling, is shown in Fig. 5. In low cycle fatigue domain (NF < 104) where the plastic part of the strain is high, the material shows shortest life at RT due to its room temperature brittleness. For lower ea, the fatigue life is shortest for specimens cycled at 800 °C. The difference between the fatigue life at 750 and 800 °C is significant especially for high cycle fatigue regime. Plotting the fatigue life data against plastic strain amplitude eap at half-life (Fig. 6) reveals a surprising effect. While at RT the usual Coffin–Manson power law is, taking into account the data scatter, valid for all plastic strain amplitudes with the exponent c = 0.25, two different regimes are observed for high temperature data. In the low cycle fatigue domain, the data can be fitted by the power law with c around 0.3. However, the curve at Fig. 6 is almost horizontal for NF > 104 for all three high testing temperatures. 3.2. Microstructure evolution in cyclic deformation
Fig. 1. Lamellar microstructure of the material. (a) TEM bright field and (b) TEM HAADF mode, white lamellas are the a2 phase.
The changes in the microstructure after fatigue failure were studied in several specimens cycled at various temperatures and strain amplitudes. It can be concluded that at RT, no changes in
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
T. Kruml, K. Obrtlík / International Journal of Fatigue xxx (2013) xxx–xxx
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Fig. 2. Cyclic hardening/softening curves. Testing temperature and strain amplitude are given in inserts.
Fig. 3. Tensile curve at RT and cyclic stress–strain curves at the four testing temperatures.
the internal structures were detected. This observation agrees with the stable cyclic response of the material. In specimens subjected to fatigue to failure at 750 °C, a microstructure appears which has never been observed in as-received state. It consists of areas of c-phase inside the grains, without lamellas of a2 phase or c-phase variants, with high dislocation density and a lot of dislocation debris in the form of small prismatic loops (Fig. 7). Formation of such structure is a consequence of cyclic plastic loading at 750 °C. The described areas are quite rare, the relative volume of destroyed lamellar structure is therefore very small. For this reason, this microstructural change is not visible on the cyclic hardening/softening curves at Fig. 2b.
Fig. 4. S–N curves at the four testing temperatures.
At 800 °C, the destruction of lamellar structure due to cyclic loading is much more frequent than at 750 °C. Contrary to the observation of specimens cycled at 750 °C, prismatic loops are not observed and dislocation density in the newly formed c phase is low (Fig. 8), substantially lower than at 750 °C (compare Figs. 7b and 8b). These c-phase islands cover approximately 10% of volume of the specimens.
4. Discussion TiAl alloys at RT usually cyclically harden [4,18]. In particular, very strong cyclic hardening was reported for alloys with 2 at.%
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
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Fig. 5. Combined fatigue life curve with guidelines.
Fig. 7. (a) Partially destroyed lamellar microstructure; and (b) high dislocation density with prismatic loops in the c phase. 750 °C, ea = 0.41%, ra = 443 MPa, foil plane parallel to the loading axis, NF = 438 cycles.
Fig. 6. Coffin–Manson diagram with guidelines.
Nb or 2Nb–2Cr [4,19]. It was explained by intensive twinning. In the present material, no cyclic hardening was measured and twins were observed only rarely. It is thus possible that the explanation of twinning induced cyclic hardening is correct and that higher amount of Nb reduces the twinning activity, e.g. by alternating the stacking fault energy. Concerning the differences of the cyclic behaviour at different testing temperatures, the changes in material response appear between 750 and 800 °C. Increase of testing temperature to 800 °C shortens the fatigue life considerably in S–N representation and material cyclically softens. This means that the maximum working temperature for applications demanding high strength of material is around to 750 °C. The change of the usual slope of Coffin–Manson diagram in low cycle fatigue domain to almost horizontal line for NF > 104 has, up to our knowledge, never been reported. On the contrary, inverse change of the slope in the Coffin–Manson diagram was observed for lamellar Ti–46.6Al–1.4Mn–2Mo cycled in air at 800 °C [20]: c = 0.22 for NF < 7 104 and c = 0.65 for NF > 7 104. Authors explained this change of the slope by change of cracking mode from transgranular to intergranular due to a2 ? c transformation on grain boundaries. In our case, intergranular cracking was not observed. However, two slopes in the Coffin–Manson diagram,
detailed for 750 °C in Fig. 9, suggests two operating mechanisms of cyclic straining whose participation depends on strain amplitude. Additional data, particularly relative to the microstructure of TiAl–8Nb alloy are necessary to explain the behaviour. The main result of the TEM microstructural observation is the documentation of changes of the microstructure of specimens cycled at 800 °C: in substantial part of the volume, lamellar microstructure is replaced by islands of c phase with low dislocation density. The destruction of lamellar microstructure occurs by gradual disappearance of a2 lamellas. As seen in Fig. 8a, in areas of 10 lm the a2 lamellas disappeared, leaving bands of dislocations on their initial places (Fig. 8b). It is supposed that with further cycling, the process would continue and resting a2 lamellas would continue to shorten, due to combined effect of temperature and cyclic straining. The disappearance of lamellar microstructure during cyclic deformation at high temperatures has already been observed by Appel et al. [11]. The microstructure of the specimen after fatigue at 850 °C published in [11] closely resembles Fig. 7a. In this work, the onset of destruction of the lamellar microstructure due to cyclic loading was assigned to 750 °C. An interesting difference between c islands formed at 750 °C and 800 °C is the different level of dislocation density including prismatic loops, which are abundant at 750 °C and non-existing at 800 °C. It can be speculated that diffusion processes, necessary for the annealing of prismatic loops, are much faster at 800 °C than at 750 °C.
Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005
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5
Stable cyclic behaviour was observed up to 750 °C, clear cyclic softening appears at 800 °C. Fatigue strength of the material is still high at 800 °C (cyclic yield stress above 300 MPa). With increasing temperature, significant reduction of fatigue life is observed in S–N diagram, particularly between 750 and 800 °C. Destruction of lamellar microstructure was observed at 750 °C and, more pronounced, at 800 °C. Areas of single c phase are formed. High density of dislocations and prismatic loops are observed at 750 °C while low dislocation density is found in these areas at 800 °C. Accentuated destruction of lamellas due to cyclic loading is responsible for cyclic softening at 800 °C.
Acknowledgements Support of the project of the Czech Science Foundation 107/11/ 0704 and by the research Project No. AV0Z 20410507 of the Academy of Sciences of the Czech Republic is acknowledged. The research was conducted in CEITEC research infrastructure supported by the project. References
Fig. 8. (a) Partially destroyed lamellar microstructure; and (b) the front of the resting lamellar microstructure and remaining dislocations in the c phase. 800 °C, ea = 0.24%, ra = 306 MPa, foil plane parallel to the loading axis, NF = 61,174 cycles.
Fig. 9. Coffin–Manson diagram for 750 °C. Dashed line is the extrapolation of data from the high cycle to low cycle domain.
5. Conclusions Low cycle fatigue behaviour of TiAl alloyed with 7.8 at.% Nb was studied from RT to 800 °C.
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Please cite this article in press as: Kruml T, Obrtlík K. Microstructure degradation in high temperature fatigue of TiAl alloy. Int J Fatigue (2013), http:// dx.doi.org/10.1016/j.ijfatigue.2013.10.005