- Email: [email protected]
Long-term creep and creep rupture characteristics of TiAl-base intermetallics
Materials Science and Engineering A 510–511 (2009) 350–355
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Long-term creep and creep rupture characteristics of TiAl-base intermetallics A. Dlouhy´ ∗ , K. Kuchaˇrová, A. Orlová Academy of Sciences CR, Institute of Physics of Materials, Zˇ iˇzkova 22, 616 62 Brno, Czech Republic
a r t i c l e
i n f o
Article history: Received 14 January 2008 Received in revised form 1 July 2008 Accepted 4 July 2008 Keywords: TiAl-base intermetallics Creep Creep rupture Monkman–Grant relationship
a b s t r a c t The present study investigates creep in three TiAl-base intermetallic alloys, Ti–48Al–2Cr–2Nb–1B, Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si and Ti–46Al–2W–0.5Si–0.7B (all at.%), in the range of rupture times up to 25 000 h. Creep behaviour of the individual alloys at 1023 K is assessed and compared and reasons for creep strength variations are discussed. Results suggest that dislocation processes contribute to the creep strain accumulation even at creep rates as low as 7 × 10−10 s−1 . Finally the applicability of Monkman–Grant-type relationships is critically analyzed. It is shown that these relationships stem from a general functional identity that is valid for any continuous creep curve. The role of the creep curve shape factor S is highlighted. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Intermetallic alloys based on the gamma TiAl compound have been developed for high temperature engineering applications [1]. In high temperature range, creep properties of the intermetallics must be taken into account since they can set limits to the performance of components. Due to their favourable strength over density ratio, these alloys are considered as a replacement for relatively heavy Ni-based superalloys in rotating engine parts [2]. Investment casting is a common processing route used for fabrication of gas turbine components. This near-net-shape technology leaves only a limited space for subsequent heat treatments to modify a coarse casting microstructure. Creep properties of cast TiAl alloy versions are therefore of primary importance. In spite of the fact that numerous investigations addressed the creep behaviour of TiAl intermetallic alloys, there has only been a limited number of studies that presented a long-term creep data [3–5]. This may be partially due to the character of data required since first, relatively short-term applications, like jet turbine TiAl blades, were considered [6]. However, with the application range extending to long-blades for stationary gas turbines [7], demand for long-term creep data has been increasing. Therefore, the present study investigates long-term creep behaviour of three different TiAl alloys based on different alloying strategies. Two investigated materials (Ti–48Al–2Cr–2Nb–1B and Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si, at.%) contain niobium while the third alloy Ti–46Al–2W–0.5Si–0.7B has been intended for long blades of the last stage in the stationary
∗ Corresponding author. Tel.: +420 532290412; fax: +420 541218657. ´ E-mail address: [email protected] (A. Dlouhy). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.07.077
turbine [8] and thus its creep strength was a prime concern. The long-term creep performance of all materials is assessed based on the well known Monkman–Grant relationship [9] and predictive capability of this relationship is discussed.
2. Experimental alloys The Ti–48Al–2Cr–2Nb–1B (at.%) alloy (alloy Nb2) was plasma melted at the Interdisciplinary Research Centre in Birmingham. The alloy was isothermally forged at the temperature 1423 K and strain rate 5 × 10−3 s−1 with a height reduction of 70%. The as-forged pancake was vacuum annealed at 1658 K for 1 h followed by furnace cooling. The resulting ␥/␣2 nearly lamellar structure is organized into colonies with the mean colony size close to 25 m and the mean lamellae width of 4.2 m. The Ti–46Al–2W–0.5Si–0.7B (at.%) alloy (alloy W) was cast by Howmet Research Corporation and supplied by ALSTOM Power Baden. The heat treatment of the as-cast cylindrical rod consisted of HIP (1533 K/172 MPa/4 h), annealing at 1623 K for 1 h and final aging at 1273 K for 6 h after which the alloy was furnace cooled to room temperature. The macrostructure in the cross-section of the heattreated rod is shown in Fig. 1. Tensile axis of the creep specimens was oriented parallel to the rod axis. Finally, the Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si (at.%) alloy (alloy Nb7) was melted, cast and delivered by Flowserve Dayton in the form of a cylindrical rod the length and diameter of which were 1200 mm and 70 mm, respectively. The rod was HIPped after casting and no additional heat treatment was applied. Tensile creep specimens were cut again parallel to the rod axis.
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
351
Fig. 1. Initial macrostructure in the cross-section of the as cast and heat-treated rod of the alloy W. Tensile creep specimens were oriented parallel with the rod axis.
Chemical compositions of the three investigated alloys and abbreviations that refer to the individual materials throughout this paper are summarized in Table 1. 3. Experimental techniques Tensile creep specimens (gauge length 25 mm and cross-section 4 mm × 3.2 mm) were cut out of the heat-treated materials using the spark erosion cutter. Constant applied stress creep tests were performed at 1023 K in a purified argon atmosphere [10,11]. Tensile elongation was continuously measured using HottingerBaldwin extensometers and corresponding strain-time readings were recorded by a PC-based data acquisition system [12]. The testing temperature was maintained constant within ±1 K along the specimen gauge length. Specimens before creep and after creep rupture were studied by light microscopy and scanning electron microscopy (SEM). 4. Results Table 2 summarizes all tensile creep data obtained in the present study for the three investigated alloys. A brief inspection of results in Table 2 shows that the longest rupture time achieved so far is slightly less than 26 000 h for the W-alloy tested at 1023 K and 150 MPa. The long-term experiment set for the Nb7-alloy at 125 MPa is still running; the accumulated exposure time at present exceeds 9800 h. Finally, the longest test taken through to rupture at 150 MPa for the Nb2-alloy ended after more than 3600 h. Our sensitive data acquisition system enables a representation of these three longest creep tests as continuous creep rate versus creep strain curves in Fig. 2. Results shown in Fig. 2 suggest that the kinetics
Fig. 2. Continuous creep curves recorded during low-stress creep of the alloys Nb2, Nb7 and W at 1023 K. The arrows indicate that the corresponding experiments are either running (alloy Nb7) or exhibit ductility exceeding the strain axis range (alloy Nb2).
of creep strain accumulation at low applied stresses is quite similar for the three investigated alloys. There always is a steep decrease of creep rate during the primary transition, which accounts for about 2% of creep strain (the instantaneous loading strain being subtracted). Beyond the point of minimum creep rate, the creep rate steadily increases and the acceleration is again comparable for all the three investigated alloys. It may be argued that the acceleration of creep after the minimum reflects contributions due to damage accumulation and premature rupture events specific for the tensile loading. However, a comparison of tensile and compression data presented in Fig. 3 clearly shows that the acceleration is mostly independent of the loading mode up to about 10% strain where final rupture event outbalances the kinetics in the tension experiment. Data in Fig. 3 were obtained for the Nb7-alloy tested at 200 MPa. We also note that even in the 25 760 h test of the W-alloy, strain accumulated to rupture exceeds 10% what is, for long-term experiments, quite a remarkable value. Double logarithmic plots of minimum creep rate versus applied stress shown in Fig. 4 suggest that, in the range of applied stresses investigated in the present study, the W-alloy exhibits superior creep strength. Changes of the minimum creep rate with the applied stress can be approximated by a power law for alloys W and Nb7. These dependences exhibit stress exponents 4.3 and 4.6 for the alloys W and Nb7, respectively. As compared to the alloy W, the alloy Nb7 creeps faster in the same range of applied stress. Diamond
Table 1 Chemical composition and nicknames of the investigated alloys. Element
Ti
Al
W
Si
Cr
Nb
Ni
C
O
N
H
B
Nb2 wt.% at.%
Bal Bal
34.3 48.6
– –
– –
2.66 1.95
4.36 1.79
– –
– –
0.064 0.150
– –
– –
0.29 1.02
Nb7 wt.% at.%
Bal Bal
29.8 45.8
– –
0.07 0.10
0.70 0.56
0.30 0.21
0.01 0.035
0.067 0.174
0.010 0.030
0.0003 0.012
– –
W wt.% at.%
Bal Bal
30.2 45.3
9.0 2.0
0.38 0.55
– –
– –
0.045 0.152
0.087 0.220
0.017 0.049
0.001 0.040
0.18 0.68
15.9 7.1 – –
352
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
Table 2 Summary of the creep and creep rupture data obtained in the present study. Alloy
Ti–48Al–2Cr–2Nb–1B (Nb2)
Ti–46Al–7Nb–0.6Cr–0.2Ni (Nb7)
Ti–46Al–2W–0.5Si–0.7B (W)
a b
Temperature, T [K]
Stress, [MPa]
Minimum creep rate, ε˙ M [s−1 ]
1023
100 150 200 300
2.6 × 10−9 1.0 × 10−8 6.0 × 10−8 1.1 × 10−6
3633.5 914.6 79.2
47.7 46.3 46.0
1023
125 175 200 250 300 350
1.6 × 10−9 1.0 × 10−8 1.0 × 10−8 2.2 × 10−8 6.9 × 10−8 2.6 × 10−7
>9800 1020.1 1974.1 1003.2 54.8 7.72
1023
150 160 180 200 220 250 250 300 300 350 360 380
6.9 × 10−10 1.2 × 10−9 2.7 × 10−9 4.9 × 10−9 6.0 × 10−9 9.0 × 10−9 1.0 × 10−8 8.0 × 10−9 1.4 × 10−8 3.4 × 10−8 4.2 × 10−8 1.0 × 10−7
25760.1 13384.3 4098.6 2600.4 2405.7 1157.8 860.5 1098.6 524.3 246.4 182.0 49.1
Time to fracture, tF [h] a
Creep fracture strain, εF [%]
Shape factor Eq. (5), SI
Shape factor Eq. (4), SD
– 0.30 0.44 0.62
0.27 0.43 0.68
6.6 13.1 14.1 2.1 1.1
– – – – – –
– 0.56 0.54 0.56 0.65 0.66
10.1 10.8 8.9 10.8 9.9 7.5 8.9 7.5 6.0 7.3 5.1 2.4
– – – – – – – – – – – –
0.63 0.54 0.45 0.42 0.52 0.50 0.35 0.42 0.44 0.41 0.54 0.74
a
b
Test interrupted in early tertiary stage. Test is still running.
data points, that represent measurements for the Nb2-alloy, can be better fitted by an exponential function, even though the slope of this function in the stress range below 200 MPa is comparable to the slopes of the remaining two power law lines. Furthermore, the isothermally forged and heat-treated nearly lamellar alloy Nb2 exhibits clearly higher creep rates than the two cast alloys Nb7 and W. This difference can exceed an order of magnitude when Nb2-alloy is compared to W-alloy. While there are distinct differences in minimum creep rates between investigated alloys (Fig. 4), stress dependences of time to fracture presented in Fig. 5 are subjected to smaller variation. Time-to-fracture data obtained for Nb-containing alloys Nb2 and Nb7 follow, within the experimental scatter, the same trend characterized by a power law with a stress exponent of −5.5. On the other
hand, the rupture times exhibited at comparable stresses by Walloy are longer even if they exhibit similar power law dependence. We note that the first data point plotted for the alloy Nb7 represents only currently accumulated time in the running creep test and an arrow indicates an expected shift of this data point for the terminal rupture state. A pronounced difference also exists as far as ductility of the investigated alloys is concerned. Rupture strains presented in Table 2 show that the nearly lamellar Nb2-alloy is fully ductile at 1023 K exhibiting rupture strains of about 45% irrespective of the applied stress level. The two cast alloys Nb7 and W fail at strains in the range 1–14% and the rupture strains seem to increase slightly with the decreasing applied stress. The Nb7-alloy was found to be prone to premature failure which causes rather high scatter of the ductility data.
Fig. 3. Comparison of strain accumulation kinetics in tension and compression for the alloy Nb7 loaded by 200 MPa at 1023 K. The compression experiment was terminated at the total strain of 39%.
Fig. 4. Stress dependences of minimum creep rate obtained at 1023 K for the alloys Nb2, Nb7 and W.
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
353
test performed in this study (see also Table 2). The double logarithmic plot shown in Fig. 6 can be satisfactorily described by Eq. (2) with the parameters mMD = 0.95 and CMD = 1.45. All experimental points fall on one master curve despite differences in chemical composition and applied stress at which individual tests were performed.
5. Discussion 5.1. Long-term creep in TiAl-based alloys
Fig. 5. Rupture times of the investigated alloys Nb2, Nb7 and W measured at 1023 K. The arrow at the first Nb7 data point indicates the running creep test and expected increase of the exposure time at rupture.
Monkman and Grant [9] were first who proposed the relationship (MGR) which links time to fracture tF and minimum creep rate ε˙ M attained in the same experiment tF (ε˙ M )
mMG
= CMG .
(1)
Later, Dobeˇs and Miliˇcka [13] argued that, in many experimental situations, the parameters CMG and mMG change with the applied stress and modified the MGR into the form tF (ε˙ M )
mMD
= CMD εF
(2)
which accounts for a possible stress dependence of the product m tF (ε˙ M ) MD due to changes of fracture strain εF with the applied stress. Miliˇcka and Dobeˇs further suggested that, in their version of the relationship (MDR), CMD and mMD are parameters independent of stress and temperature [14]. Fig. 6 presents a relation between mean creep rate εF /tF and the minimum creep rate ε˙ M of each creep
Fig. 6. Relation between mean creep rate εF /tF and the minimum creep rate ε˙ M demonstrates the general validity of Eq. (2) for each creep test taken through to rupture in this study.
A remarkable feature of the longest creep test reported in the present study (the W alloy crept at 1023 K under 150 MPa, time to fracture 25 760 h) is a relatively high fracture strain which exceeds 10%. Such a high ductility cannot be an outcome of diffusion creep only. Dislocation processes of whatever nature must contribute during the accumulation of this strain. This view is supported by the fact that the corresponding dependence of minimum creep rate on the applied stress in Fig. 4 is straight in double logarithmic co-ordinates. The plot exhibits the stress exponent close to 4.5 throughout the investigated stress range and there is no indication of any transition to a less steep dependence in the region of low applied stresses. We note that lower stress exponents approaching unity are typical for diffusion creep [15]. A pronounced deceleration of creep rate during the primary transition, shown in Fig. 2, is typical not only for the W-alloy but also for the other two alloys Nb7 and Nb2 tested at low applied stresses. Primary strains needed to reach the minimum creep rate were 2.2, 2.1 and 1.3% for the alloys Nb2, Nb7 and W, respectively. The creep rate falls by at least three orders of magnitude within these strain intervals. Therefore, the underlying dislocation process must also slow down considerably. In view of the limited primary strains, this process cannot be most probably associated with the classical formation of dislocation substructure that would require much larger deformations. Instead, a lack of superdislocations and insufficient twinning activity that, if present, both contribute to plastic displacements parallel with the c-axis of the L10 gamma-TiAl lattice, may cause the formation of incompatibility stresses and deceleration of primary creep [16]. As it is documented in Fig. 2, the largest portion of total strains measured in present low-stress experiments accumulates in the stage where creep rate slowly but steadily increases. It has been shown (Fig. 3) that the observed acceleration is largely independent of a loading mode (tension or compression) and cannot thus be fully accounted for on the basis of damage processes, like cavitation, that would be specific for tensile creep only. We also note that the character (kinetics) of strain accumulation during tertiary creep illustrated in Fig. 2 is very similar for all the three alloys. Unless this is a pure coincidence, the creep rate acceleration throughout the tertiary stage does not seem to be much influenced by the particular alloy chemistry and a specific alloy microstructure [17]. Therefore, some more or less common dislocation mechanism should govern creep in the investigated intermetallics also in the tertiary stage. The evolution of deformation structures during creep in gamma TiAl based intermetallics was recently a subject of systematic quantitative studies. It has been shown that similar acceleration of creep rate during tertiary stage can be accounted for on the basis of increasing activity of superdislocations and/or deformation twins [18,19]. Whether these microstructural mechanisms are also relevant for low-stress long-term creep in TiAl based alloys is at present an open question [20]. Further quantitative microstructural studies are needed to clarify this point. As it is documented in Fig. 4, the alloy chemistry and heat treatment clearly influence creep strength of the investigated alloys. Recent investigations confirmed that there are some common features but also relevant differences in the alloy microstructures [17].
354
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
where
1
y(x) dx.
S(y) =
(5)
0
Fig. 7. Rupture surface observed after long-term exposure of the alloy W (crept at 1023 K under 150 MPa, time to fracture 25 760 h). Numerous {1 1 1}-␥ and {0 0 0 1}␣2 facets are clearly visible in the SEM micrograph.
Even though dislocation processes operating during creep of the investigated materials may be quite similar as it has been discussed above, the resistance to these processes increases from the fast creeping alloy Nb2 towards the rather strong alloy W. There always are lamellar ␥/␣2 grains in all the three intermetallics that are mixed with the ␥ grains in the duplex microstructures of the cast alloys Nb7 and W [17]. The superior creep resistance of the alloy W is certainly associated with a presence of silicide and W-rich B2 particles. These remarkably stable particles precipitate in the gamma matrix but also decorate lamellar ␥/␣2 interfaces. Consequently, these particles limit the motion of matrix and interface dislocations [17,21]. Lamellar interfaces also play an important role in rupture of longterm crept specimens. {1 1 1}-␥ and {0 0 0 1}-␣2 facets are clearly visible in Fig. 7, which illustrates the rupture surface of the alloy W after exposure lasting for 25 760 h (see Table 2 for more details). Further work is needed to (i) rationalize the role of lamellar interfaces during creep rupture of TiAl-based alloys and (ii) account for differences in creep rates between the two Nb-containing alloys. 5.2. Applicability of Monkman–Grant relationship There has traditionally been a quest for a relationship that would allow a successful prediction of time to fracture tF in long-term creep tests from shorter term creep data. The minimum creep rate establishes early in creep life in long-term tests, therefore, the prediction of tF based on the minimum creep rate (Eq. (1)) would save expensive testing time. This prediction strategy often fails since parameters CMG and mMG do generally change with the applied stress [13]. However, the modification suggested by Dobeˇs and Miliˇcka [13,14] does not improve the prediction capability of the relationship since, instead of one long-term parameter – the time to fracture tF , their relationship requires also the second long-term parameter – the fracture strain εF . Here we show that both MGR and MDR are a direct consequence of a functional identity which is fulfilled by every smooth creep curve. When a non-zero positive creep rate ε˙ = dε/dt can be calculated in any point of a creep curve ε(t), then time to fracture tF is
tF =
tF
dt = 0
0
εF
dε . ˙ ε(ε)
6. Summary and conclusions Creep characteristics of Ti–48Al–2Cr–2Nb–1B (Nb2), Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si (Nb7) and Ti–46Al–2W–0.5Si– 0.7B (W) (all at.%) intermetallic alloys have been investigated at 1023 K including long-term exposures in which rupture time extended beyond 25 000 h. The performance of the three materials has been compared and the applicability of the Monkman–Granttype relationships for the assessment and prediction of long-term creep data has been discussed. Results obtained in the present study can be summarized as follows:
(3)
˙ this identity Using normalized co-ordinates x = ε/εF and y = ε˙ M /ε, can be transformed into the form tF ε˙ M = S(y) εF
The values of the functional S(y) fall into the interval (0, 1) and depend on the “shape” of the normalized curve y(x). Now, when mMD = 1 and CMD = S, the identity Eq. (4) casts into the MDR (Eq. (2)). Furthermore, with mMG = 1 and CMG = SεF the MGR (Eq. (1)) is obtained. In summary, it is evident that both relationships (Eq. (1) and Eq. (2)) stem from the general expression given in Eq. (4). This association has several important consequences. As it has been reported in experimental studies, e.g. [14], the values of mMG and mMD are often different from unity. One particular example is given in this study, where fitting of data in Fig. 6 yields mMD = 0.95. The difference of these exponents from unity compensates for the fact that CMG and CMD are (incorrectly) taken as constants in MGR and MDR. Even the functional S, which has been calculated for the alloy Nb2 using Eq. (5), does change with the applied stress. Results of these calculations are presented in Table 2 in the SI column. Equivalently, shape factors S can also be obtained using the formula given by the left hand side of Eq. (4). Table 2 also summarizes these values for all the three investigated alloys in the SD column. There is no doubt that the MGR (Eq. (1)) has its practical value as it was originally proposed by Monkman and Grant [9]. However, its predictive potential depends very much upon how reliably we can extrapolate the stress dependence of the strain to fracture εF and the shape factor S from the high stress range (short term creep test) down to the range of low applied stresses. This extrapolation may represent a difficulty as it has been demonstrated in this study. Every creep curve results from a particular combination of microstructural processes that govern creep deformation, damage accumulation and creep rupture. Whatever this combination is, the corresponding creep curve must obey the identity relation Eq. (4) and, in certain cases, also its special forms given by either MGR or MDR with mMG = mMD = 1. The nature and mutual interplay of deformation and rupture modes thus is not important for the validity of Eq. (4) and subsequently also for the validity of Eqs. (1) and (2). Therefore, no deeper insight into the governing processes can be obtained solely on the basis of MGR and MDR. Careful microstructural investigations are always needed to identify the relevant microstructural mechanisms. As early as in 1985 Povolo [22], using a similar line of arguments as the ones put forward in Eqs. (3)–(5), criticized attempts aimed at a microstructurally based explanation of the MGR and MDR. It seems that many following authors have not taken Povolo’s fair analysis into account.
(4)
(i) Creep strength of the investigated materials, characterized by means of the inverse minimum creep rate, increases in the succession Nb2, Nb7 and W. The superior creep strength of the W-alloy is associated with the presence of stable silicide and B2 particles [17].
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
(ii) Rupture strain accumulated during the long-term creep of the alloy W exceeded 10%. This result suggests that dislocation mechanisms are still important at conditions that result in rupture times exceeding 25 000 h. (iii) Creep strain accumulation kinetics is similar in all three materials at comparable stresses and strains. Therefore similar dislocation mechanisms are expected to govern creep of these alloys. (iv) Rupture surfaces observed after long-term creep exhibit numerous {1 1 1}-␥ and {0 0 0 1}-␣2 facets. Rupture along lamellar interfaces and associated microstructural mechanisms should be a subject of further investigations. (v) Monkman–Grant and Miliˇcka–Dobeˇs relationships are a direct consequence of the functional identity described by Eq. (4). These relationships should be used with caution because their predictive ability depends on the changes of rupture strain and the creep curve shape factor S with the applied stress.
355
References [1] [2] [3] [4] [5] [6] [7] [8]
[9] [10] [11] [12] [13] [14] [15] [16]
Acknowledgements
[17] [18]
Dr. M. Nazmy, Alstom-Power Switzerland, kindly provided the W alloy. The authors also thank Dr. M. Svoboda who assisted investigations of rupture surfaces using SEM. The financial support was received from the Czech Science Foundation under the contract no. 106/07/0762.
[20] [21]
[19]
[22]
Y.-W. Kim, Intermetallics 6 (1998) 623–628. M. Nazmy, M. Staubli, US Pat. no. 5,207,982 and European Pat. no. 4555005 B1. F. Appel, Intermetallics 9 (2001) 907–914. J. Lapin, M. Nazmy, Mater. Sci. Eng. A380 (2004) 298–307. J. Beddoes, D.Y. Seo, H. Saari, Scripta Mater. 52 (2005) 745–750. C.M. Austin, Curr. Opin. Solid State Mater. Sci. 4 (1999) 239–242. M. Nazmy, M. Staubli, G. Onofrio, V. Lupinc, Scripta Mater. 45 (2001) 787– 792. ´ D. Arrell, B. Karlsson, J. Lapin, V. Lupinc, M. Nazmy, K. Nikbin, A. Dlouhy, M. Staubli, in: J. Lecomte-Beckers, et al. (Eds.), Materials for Advanced Power Engineering 2002, Liège, Belgium, September 30–October 2, 2002, Forschungszentrum Jülich, Jülich, Germany, 2002, pp. 605–614. F.C. Monkman, N.J. Grant, Proc. ASTM 56 (1956) 593–605. ˇ ´ J. Cadek, T. Hostinsky, J. Test. Eval. 4 (1976) 26–28. ˇ F. Dobeˇs, O. Zvˇerˇina, J. Cadek, Metallic Mater. 24 (1986) 599–610. K. Miliˇcka, Internal Report, IPM AS CR, Brno, Czech Republic, 1993. F. Dobeˇs, K. Miliˇcka, Met. Sci. 10 (1976) 382–384. K. Miliˇcka, F. Dobeˇs, Metallic Mater. 15 (1977) 186–200. ˇ J. Cadek, Creep in Metallic Materials, Elsevier, Amsterdam, Netherlands, 1988. ´ K. Kuchaˇrová, J. Bˇrezina, Mater. Sci. Eng. A319–321 (2001) 820– A. Dlouhy, 826. ´ K. Kuchaˇrová, Intermetallics 12 (2004) 705–711. A. Dlouhy, ´ Zeitschrift für Metallkd. 95 (2004) A. Orlová, K. Kuchaˇrová, A. Dlouhy, 1011–1019. ´ A. Orlová, K. Kuchaˇrová, A. Dlouhy, Mater. Sci. Eng. A483–484 (2008) 109– 112. J. Lapin, Intermetallics 14 (2006) 115–122. L.M. Hsiung, T.G. Nieh, B.W. Choi, J. Wadsworth, Mater. Sci. Eng. A329–331 (2002) 637–643. F. Povolo, J. Mater. Sci. 20 (1985) 2005–2010.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Long-term creep and creep rupture characteristics of TiAl-base intermetallics A. Dlouhy´ ∗ , K. Kuchaˇrová, A. Orlová Academy of Sciences CR, Institute of Physics of Materials, Zˇ iˇzkova 22, 616 62 Brno, Czech Republic
a r t i c l e
i n f o
Article history: Received 14 January 2008 Received in revised form 1 July 2008 Accepted 4 July 2008 Keywords: TiAl-base intermetallics Creep Creep rupture Monkman–Grant relationship
a b s t r a c t The present study investigates creep in three TiAl-base intermetallic alloys, Ti–48Al–2Cr–2Nb–1B, Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si and Ti–46Al–2W–0.5Si–0.7B (all at.%), in the range of rupture times up to 25 000 h. Creep behaviour of the individual alloys at 1023 K is assessed and compared and reasons for creep strength variations are discussed. Results suggest that dislocation processes contribute to the creep strain accumulation even at creep rates as low as 7 × 10−10 s−1 . Finally the applicability of Monkman–Grant-type relationships is critically analyzed. It is shown that these relationships stem from a general functional identity that is valid for any continuous creep curve. The role of the creep curve shape factor S is highlighted. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Intermetallic alloys based on the gamma TiAl compound have been developed for high temperature engineering applications [1]. In high temperature range, creep properties of the intermetallics must be taken into account since they can set limits to the performance of components. Due to their favourable strength over density ratio, these alloys are considered as a replacement for relatively heavy Ni-based superalloys in rotating engine parts [2]. Investment casting is a common processing route used for fabrication of gas turbine components. This near-net-shape technology leaves only a limited space for subsequent heat treatments to modify a coarse casting microstructure. Creep properties of cast TiAl alloy versions are therefore of primary importance. In spite of the fact that numerous investigations addressed the creep behaviour of TiAl intermetallic alloys, there has only been a limited number of studies that presented a long-term creep data [3–5]. This may be partially due to the character of data required since first, relatively short-term applications, like jet turbine TiAl blades, were considered [6]. However, with the application range extending to long-blades for stationary gas turbines [7], demand for long-term creep data has been increasing. Therefore, the present study investigates long-term creep behaviour of three different TiAl alloys based on different alloying strategies. Two investigated materials (Ti–48Al–2Cr–2Nb–1B and Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si, at.%) contain niobium while the third alloy Ti–46Al–2W–0.5Si–0.7B has been intended for long blades of the last stage in the stationary
∗ Corresponding author. Tel.: +420 532290412; fax: +420 541218657. ´ E-mail address: [email protected] (A. Dlouhy). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.07.077
turbine [8] and thus its creep strength was a prime concern. The long-term creep performance of all materials is assessed based on the well known Monkman–Grant relationship [9] and predictive capability of this relationship is discussed.
2. Experimental alloys The Ti–48Al–2Cr–2Nb–1B (at.%) alloy (alloy Nb2) was plasma melted at the Interdisciplinary Research Centre in Birmingham. The alloy was isothermally forged at the temperature 1423 K and strain rate 5 × 10−3 s−1 with a height reduction of 70%. The as-forged pancake was vacuum annealed at 1658 K for 1 h followed by furnace cooling. The resulting ␥/␣2 nearly lamellar structure is organized into colonies with the mean colony size close to 25 m and the mean lamellae width of 4.2 m. The Ti–46Al–2W–0.5Si–0.7B (at.%) alloy (alloy W) was cast by Howmet Research Corporation and supplied by ALSTOM Power Baden. The heat treatment of the as-cast cylindrical rod consisted of HIP (1533 K/172 MPa/4 h), annealing at 1623 K for 1 h and final aging at 1273 K for 6 h after which the alloy was furnace cooled to room temperature. The macrostructure in the cross-section of the heattreated rod is shown in Fig. 1. Tensile axis of the creep specimens was oriented parallel to the rod axis. Finally, the Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si (at.%) alloy (alloy Nb7) was melted, cast and delivered by Flowserve Dayton in the form of a cylindrical rod the length and diameter of which were 1200 mm and 70 mm, respectively. The rod was HIPped after casting and no additional heat treatment was applied. Tensile creep specimens were cut again parallel to the rod axis.
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
351
Fig. 1. Initial macrostructure in the cross-section of the as cast and heat-treated rod of the alloy W. Tensile creep specimens were oriented parallel with the rod axis.
Chemical compositions of the three investigated alloys and abbreviations that refer to the individual materials throughout this paper are summarized in Table 1. 3. Experimental techniques Tensile creep specimens (gauge length 25 mm and cross-section 4 mm × 3.2 mm) were cut out of the heat-treated materials using the spark erosion cutter. Constant applied stress creep tests were performed at 1023 K in a purified argon atmosphere [10,11]. Tensile elongation was continuously measured using HottingerBaldwin extensometers and corresponding strain-time readings were recorded by a PC-based data acquisition system [12]. The testing temperature was maintained constant within ±1 K along the specimen gauge length. Specimens before creep and after creep rupture were studied by light microscopy and scanning electron microscopy (SEM). 4. Results Table 2 summarizes all tensile creep data obtained in the present study for the three investigated alloys. A brief inspection of results in Table 2 shows that the longest rupture time achieved so far is slightly less than 26 000 h for the W-alloy tested at 1023 K and 150 MPa. The long-term experiment set for the Nb7-alloy at 125 MPa is still running; the accumulated exposure time at present exceeds 9800 h. Finally, the longest test taken through to rupture at 150 MPa for the Nb2-alloy ended after more than 3600 h. Our sensitive data acquisition system enables a representation of these three longest creep tests as continuous creep rate versus creep strain curves in Fig. 2. Results shown in Fig. 2 suggest that the kinetics
Fig. 2. Continuous creep curves recorded during low-stress creep of the alloys Nb2, Nb7 and W at 1023 K. The arrows indicate that the corresponding experiments are either running (alloy Nb7) or exhibit ductility exceeding the strain axis range (alloy Nb2).
of creep strain accumulation at low applied stresses is quite similar for the three investigated alloys. There always is a steep decrease of creep rate during the primary transition, which accounts for about 2% of creep strain (the instantaneous loading strain being subtracted). Beyond the point of minimum creep rate, the creep rate steadily increases and the acceleration is again comparable for all the three investigated alloys. It may be argued that the acceleration of creep after the minimum reflects contributions due to damage accumulation and premature rupture events specific for the tensile loading. However, a comparison of tensile and compression data presented in Fig. 3 clearly shows that the acceleration is mostly independent of the loading mode up to about 10% strain where final rupture event outbalances the kinetics in the tension experiment. Data in Fig. 3 were obtained for the Nb7-alloy tested at 200 MPa. We also note that even in the 25 760 h test of the W-alloy, strain accumulated to rupture exceeds 10% what is, for long-term experiments, quite a remarkable value. Double logarithmic plots of minimum creep rate versus applied stress shown in Fig. 4 suggest that, in the range of applied stresses investigated in the present study, the W-alloy exhibits superior creep strength. Changes of the minimum creep rate with the applied stress can be approximated by a power law for alloys W and Nb7. These dependences exhibit stress exponents 4.3 and 4.6 for the alloys W and Nb7, respectively. As compared to the alloy W, the alloy Nb7 creeps faster in the same range of applied stress. Diamond
Table 1 Chemical composition and nicknames of the investigated alloys. Element
Ti
Al
W
Si
Cr
Nb
Ni
C
O
N
H
B
Nb2 wt.% at.%
Bal Bal
34.3 48.6
– –
– –
2.66 1.95
4.36 1.79
– –
– –
0.064 0.150
– –
– –
0.29 1.02
Nb7 wt.% at.%
Bal Bal
29.8 45.8
– –
0.07 0.10
0.70 0.56
0.30 0.21
0.01 0.035
0.067 0.174
0.010 0.030
0.0003 0.012
– –
W wt.% at.%
Bal Bal
30.2 45.3
9.0 2.0
0.38 0.55
– –
– –
0.045 0.152
0.087 0.220
0.017 0.049
0.001 0.040
0.18 0.68
15.9 7.1 – –
352
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
Table 2 Summary of the creep and creep rupture data obtained in the present study. Alloy
Ti–48Al–2Cr–2Nb–1B (Nb2)
Ti–46Al–7Nb–0.6Cr–0.2Ni (Nb7)
Ti–46Al–2W–0.5Si–0.7B (W)
a b
Temperature, T [K]
Stress, [MPa]
Minimum creep rate, ε˙ M [s−1 ]
1023
100 150 200 300
2.6 × 10−9 1.0 × 10−8 6.0 × 10−8 1.1 × 10−6
3633.5 914.6 79.2
47.7 46.3 46.0
1023
125 175 200 250 300 350
1.6 × 10−9 1.0 × 10−8 1.0 × 10−8 2.2 × 10−8 6.9 × 10−8 2.6 × 10−7
>9800 1020.1 1974.1 1003.2 54.8 7.72
1023
150 160 180 200 220 250 250 300 300 350 360 380
6.9 × 10−10 1.2 × 10−9 2.7 × 10−9 4.9 × 10−9 6.0 × 10−9 9.0 × 10−9 1.0 × 10−8 8.0 × 10−9 1.4 × 10−8 3.4 × 10−8 4.2 × 10−8 1.0 × 10−7
25760.1 13384.3 4098.6 2600.4 2405.7 1157.8 860.5 1098.6 524.3 246.4 182.0 49.1
Time to fracture, tF [h] a
Creep fracture strain, εF [%]
Shape factor Eq. (5), SI
Shape factor Eq. (4), SD
– 0.30 0.44 0.62
0.27 0.43 0.68
6.6 13.1 14.1 2.1 1.1
– – – – – –
– 0.56 0.54 0.56 0.65 0.66
10.1 10.8 8.9 10.8 9.9 7.5 8.9 7.5 6.0 7.3 5.1 2.4
– – – – – – – – – – – –
0.63 0.54 0.45 0.42 0.52 0.50 0.35 0.42 0.44 0.41 0.54 0.74
a
b
Test interrupted in early tertiary stage. Test is still running.
data points, that represent measurements for the Nb2-alloy, can be better fitted by an exponential function, even though the slope of this function in the stress range below 200 MPa is comparable to the slopes of the remaining two power law lines. Furthermore, the isothermally forged and heat-treated nearly lamellar alloy Nb2 exhibits clearly higher creep rates than the two cast alloys Nb7 and W. This difference can exceed an order of magnitude when Nb2-alloy is compared to W-alloy. While there are distinct differences in minimum creep rates between investigated alloys (Fig. 4), stress dependences of time to fracture presented in Fig. 5 are subjected to smaller variation. Time-to-fracture data obtained for Nb-containing alloys Nb2 and Nb7 follow, within the experimental scatter, the same trend characterized by a power law with a stress exponent of −5.5. On the other
hand, the rupture times exhibited at comparable stresses by Walloy are longer even if they exhibit similar power law dependence. We note that the first data point plotted for the alloy Nb7 represents only currently accumulated time in the running creep test and an arrow indicates an expected shift of this data point for the terminal rupture state. A pronounced difference also exists as far as ductility of the investigated alloys is concerned. Rupture strains presented in Table 2 show that the nearly lamellar Nb2-alloy is fully ductile at 1023 K exhibiting rupture strains of about 45% irrespective of the applied stress level. The two cast alloys Nb7 and W fail at strains in the range 1–14% and the rupture strains seem to increase slightly with the decreasing applied stress. The Nb7-alloy was found to be prone to premature failure which causes rather high scatter of the ductility data.
Fig. 3. Comparison of strain accumulation kinetics in tension and compression for the alloy Nb7 loaded by 200 MPa at 1023 K. The compression experiment was terminated at the total strain of 39%.
Fig. 4. Stress dependences of minimum creep rate obtained at 1023 K for the alloys Nb2, Nb7 and W.
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
353
test performed in this study (see also Table 2). The double logarithmic plot shown in Fig. 6 can be satisfactorily described by Eq. (2) with the parameters mMD = 0.95 and CMD = 1.45. All experimental points fall on one master curve despite differences in chemical composition and applied stress at which individual tests were performed.
5. Discussion 5.1. Long-term creep in TiAl-based alloys
Fig. 5. Rupture times of the investigated alloys Nb2, Nb7 and W measured at 1023 K. The arrow at the first Nb7 data point indicates the running creep test and expected increase of the exposure time at rupture.
Monkman and Grant [9] were first who proposed the relationship (MGR) which links time to fracture tF and minimum creep rate ε˙ M attained in the same experiment tF (ε˙ M )
mMG
= CMG .
(1)
Later, Dobeˇs and Miliˇcka [13] argued that, in many experimental situations, the parameters CMG and mMG change with the applied stress and modified the MGR into the form tF (ε˙ M )
mMD
= CMD εF
(2)
which accounts for a possible stress dependence of the product m tF (ε˙ M ) MD due to changes of fracture strain εF with the applied stress. Miliˇcka and Dobeˇs further suggested that, in their version of the relationship (MDR), CMD and mMD are parameters independent of stress and temperature [14]. Fig. 6 presents a relation between mean creep rate εF /tF and the minimum creep rate ε˙ M of each creep
Fig. 6. Relation between mean creep rate εF /tF and the minimum creep rate ε˙ M demonstrates the general validity of Eq. (2) for each creep test taken through to rupture in this study.
A remarkable feature of the longest creep test reported in the present study (the W alloy crept at 1023 K under 150 MPa, time to fracture 25 760 h) is a relatively high fracture strain which exceeds 10%. Such a high ductility cannot be an outcome of diffusion creep only. Dislocation processes of whatever nature must contribute during the accumulation of this strain. This view is supported by the fact that the corresponding dependence of minimum creep rate on the applied stress in Fig. 4 is straight in double logarithmic co-ordinates. The plot exhibits the stress exponent close to 4.5 throughout the investigated stress range and there is no indication of any transition to a less steep dependence in the region of low applied stresses. We note that lower stress exponents approaching unity are typical for diffusion creep [15]. A pronounced deceleration of creep rate during the primary transition, shown in Fig. 2, is typical not only for the W-alloy but also for the other two alloys Nb7 and Nb2 tested at low applied stresses. Primary strains needed to reach the minimum creep rate were 2.2, 2.1 and 1.3% for the alloys Nb2, Nb7 and W, respectively. The creep rate falls by at least three orders of magnitude within these strain intervals. Therefore, the underlying dislocation process must also slow down considerably. In view of the limited primary strains, this process cannot be most probably associated with the classical formation of dislocation substructure that would require much larger deformations. Instead, a lack of superdislocations and insufficient twinning activity that, if present, both contribute to plastic displacements parallel with the c-axis of the L10 gamma-TiAl lattice, may cause the formation of incompatibility stresses and deceleration of primary creep [16]. As it is documented in Fig. 2, the largest portion of total strains measured in present low-stress experiments accumulates in the stage where creep rate slowly but steadily increases. It has been shown (Fig. 3) that the observed acceleration is largely independent of a loading mode (tension or compression) and cannot thus be fully accounted for on the basis of damage processes, like cavitation, that would be specific for tensile creep only. We also note that the character (kinetics) of strain accumulation during tertiary creep illustrated in Fig. 2 is very similar for all the three alloys. Unless this is a pure coincidence, the creep rate acceleration throughout the tertiary stage does not seem to be much influenced by the particular alloy chemistry and a specific alloy microstructure [17]. Therefore, some more or less common dislocation mechanism should govern creep in the investigated intermetallics also in the tertiary stage. The evolution of deformation structures during creep in gamma TiAl based intermetallics was recently a subject of systematic quantitative studies. It has been shown that similar acceleration of creep rate during tertiary stage can be accounted for on the basis of increasing activity of superdislocations and/or deformation twins [18,19]. Whether these microstructural mechanisms are also relevant for low-stress long-term creep in TiAl based alloys is at present an open question [20]. Further quantitative microstructural studies are needed to clarify this point. As it is documented in Fig. 4, the alloy chemistry and heat treatment clearly influence creep strength of the investigated alloys. Recent investigations confirmed that there are some common features but also relevant differences in the alloy microstructures [17].
354
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
where
1
y(x) dx.
S(y) =
(5)
0
Fig. 7. Rupture surface observed after long-term exposure of the alloy W (crept at 1023 K under 150 MPa, time to fracture 25 760 h). Numerous {1 1 1}-␥ and {0 0 0 1}␣2 facets are clearly visible in the SEM micrograph.
Even though dislocation processes operating during creep of the investigated materials may be quite similar as it has been discussed above, the resistance to these processes increases from the fast creeping alloy Nb2 towards the rather strong alloy W. There always are lamellar ␥/␣2 grains in all the three intermetallics that are mixed with the ␥ grains in the duplex microstructures of the cast alloys Nb7 and W [17]. The superior creep resistance of the alloy W is certainly associated with a presence of silicide and W-rich B2 particles. These remarkably stable particles precipitate in the gamma matrix but also decorate lamellar ␥/␣2 interfaces. Consequently, these particles limit the motion of matrix and interface dislocations [17,21]. Lamellar interfaces also play an important role in rupture of longterm crept specimens. {1 1 1}-␥ and {0 0 0 1}-␣2 facets are clearly visible in Fig. 7, which illustrates the rupture surface of the alloy W after exposure lasting for 25 760 h (see Table 2 for more details). Further work is needed to (i) rationalize the role of lamellar interfaces during creep rupture of TiAl-based alloys and (ii) account for differences in creep rates between the two Nb-containing alloys. 5.2. Applicability of Monkman–Grant relationship There has traditionally been a quest for a relationship that would allow a successful prediction of time to fracture tF in long-term creep tests from shorter term creep data. The minimum creep rate establishes early in creep life in long-term tests, therefore, the prediction of tF based on the minimum creep rate (Eq. (1)) would save expensive testing time. This prediction strategy often fails since parameters CMG and mMG do generally change with the applied stress [13]. However, the modification suggested by Dobeˇs and Miliˇcka [13,14] does not improve the prediction capability of the relationship since, instead of one long-term parameter – the time to fracture tF , their relationship requires also the second long-term parameter – the fracture strain εF . Here we show that both MGR and MDR are a direct consequence of a functional identity which is fulfilled by every smooth creep curve. When a non-zero positive creep rate ε˙ = dε/dt can be calculated in any point of a creep curve ε(t), then time to fracture tF is
tF =
tF
dt = 0
0
εF
dε . ˙ ε(ε)
6. Summary and conclusions Creep characteristics of Ti–48Al–2Cr–2Nb–1B (Nb2), Ti–46Al–7Nb–0.6Cr–0.2Ni–0.1Si (Nb7) and Ti–46Al–2W–0.5Si– 0.7B (W) (all at.%) intermetallic alloys have been investigated at 1023 K including long-term exposures in which rupture time extended beyond 25 000 h. The performance of the three materials has been compared and the applicability of the Monkman–Granttype relationships for the assessment and prediction of long-term creep data has been discussed. Results obtained in the present study can be summarized as follows:
(3)
˙ this identity Using normalized co-ordinates x = ε/εF and y = ε˙ M /ε, can be transformed into the form tF ε˙ M = S(y) εF
The values of the functional S(y) fall into the interval (0, 1) and depend on the “shape” of the normalized curve y(x). Now, when mMD = 1 and CMD = S, the identity Eq. (4) casts into the MDR (Eq. (2)). Furthermore, with mMG = 1 and CMG = SεF the MGR (Eq. (1)) is obtained. In summary, it is evident that both relationships (Eq. (1) and Eq. (2)) stem from the general expression given in Eq. (4). This association has several important consequences. As it has been reported in experimental studies, e.g. [14], the values of mMG and mMD are often different from unity. One particular example is given in this study, where fitting of data in Fig. 6 yields mMD = 0.95. The difference of these exponents from unity compensates for the fact that CMG and CMD are (incorrectly) taken as constants in MGR and MDR. Even the functional S, which has been calculated for the alloy Nb2 using Eq. (5), does change with the applied stress. Results of these calculations are presented in Table 2 in the SI column. Equivalently, shape factors S can also be obtained using the formula given by the left hand side of Eq. (4). Table 2 also summarizes these values for all the three investigated alloys in the SD column. There is no doubt that the MGR (Eq. (1)) has its practical value as it was originally proposed by Monkman and Grant [9]. However, its predictive potential depends very much upon how reliably we can extrapolate the stress dependence of the strain to fracture εF and the shape factor S from the high stress range (short term creep test) down to the range of low applied stresses. This extrapolation may represent a difficulty as it has been demonstrated in this study. Every creep curve results from a particular combination of microstructural processes that govern creep deformation, damage accumulation and creep rupture. Whatever this combination is, the corresponding creep curve must obey the identity relation Eq. (4) and, in certain cases, also its special forms given by either MGR or MDR with mMG = mMD = 1. The nature and mutual interplay of deformation and rupture modes thus is not important for the validity of Eq. (4) and subsequently also for the validity of Eqs. (1) and (2). Therefore, no deeper insight into the governing processes can be obtained solely on the basis of MGR and MDR. Careful microstructural investigations are always needed to identify the relevant microstructural mechanisms. As early as in 1985 Povolo [22], using a similar line of arguments as the ones put forward in Eqs. (3)–(5), criticized attempts aimed at a microstructurally based explanation of the MGR and MDR. It seems that many following authors have not taken Povolo’s fair analysis into account.
(4)
(i) Creep strength of the investigated materials, characterized by means of the inverse minimum creep rate, increases in the succession Nb2, Nb7 and W. The superior creep strength of the W-alloy is associated with the presence of stable silicide and B2 particles [17].
A. Dlouh´ y et al. / Materials Science and Engineering A 510–511 (2009) 350–355
(ii) Rupture strain accumulated during the long-term creep of the alloy W exceeded 10%. This result suggests that dislocation mechanisms are still important at conditions that result in rupture times exceeding 25 000 h. (iii) Creep strain accumulation kinetics is similar in all three materials at comparable stresses and strains. Therefore similar dislocation mechanisms are expected to govern creep of these alloys. (iv) Rupture surfaces observed after long-term creep exhibit numerous {1 1 1}-␥ and {0 0 0 1}-␣2 facets. Rupture along lamellar interfaces and associated microstructural mechanisms should be a subject of further investigations. (v) Monkman–Grant and Miliˇcka–Dobeˇs relationships are a direct consequence of the functional identity described by Eq. (4). These relationships should be used with caution because their predictive ability depends on the changes of rupture strain and the creep curve shape factor S with the applied stress.
355
References [1] [2] [3] [4] [5] [6] [7] [8]
[9] [10] [11] [12] [13] [14] [15] [16]
Acknowledgements
[17] [18]
Dr. M. Nazmy, Alstom-Power Switzerland, kindly provided the W alloy. The authors also thank Dr. M. Svoboda who assisted investigations of rupture surfaces using SEM. The financial support was received from the Czech Science Foundation under the contract no. 106/07/0762.
[20] [21]
[19]
[22]
Y.-W. Kim, Intermetallics 6 (1998) 623–628. M. Nazmy, M. Staubli, US Pat. no. 5,207,982 and European Pat. no. 4555005 B1. F. Appel, Intermetallics 9 (2001) 907–914. J. Lapin, M. Nazmy, Mater. Sci. Eng. A380 (2004) 298–307. J. Beddoes, D.Y. Seo, H. Saari, Scripta Mater. 52 (2005) 745–750. C.M. Austin, Curr. Opin. Solid State Mater. Sci. 4 (1999) 239–242. M. Nazmy, M. Staubli, G. Onofrio, V. Lupinc, Scripta Mater. 45 (2001) 787– 792. ´ D. Arrell, B. Karlsson, J. Lapin, V. Lupinc, M. Nazmy, K. Nikbin, A. Dlouhy, M. Staubli, in: J. Lecomte-Beckers, et al. (Eds.), Materials for Advanced Power Engineering 2002, Liège, Belgium, September 30–October 2, 2002, Forschungszentrum Jülich, Jülich, Germany, 2002, pp. 605–614. F.C. Monkman, N.J. Grant, Proc. ASTM 56 (1956) 593–605. ˇ ´ J. Cadek, T. Hostinsky, J. Test. Eval. 4 (1976) 26–28. ˇ F. Dobeˇs, O. Zvˇerˇina, J. Cadek, Metallic Mater. 24 (1986) 599–610. K. Miliˇcka, Internal Report, IPM AS CR, Brno, Czech Republic, 1993. F. Dobeˇs, K. Miliˇcka, Met. Sci. 10 (1976) 382–384. K. Miliˇcka, F. Dobeˇs, Metallic Mater. 15 (1977) 186–200. ˇ J. Cadek, Creep in Metallic Materials, Elsevier, Amsterdam, Netherlands, 1988. ´ K. Kuchaˇrová, J. Bˇrezina, Mater. Sci. Eng. A319–321 (2001) 820– A. Dlouhy, 826. ´ K. Kuchaˇrová, Intermetallics 12 (2004) 705–711. A. Dlouhy, ´ Zeitschrift für Metallkd. 95 (2004) A. Orlová, K. Kuchaˇrová, A. Dlouhy, 1011–1019. ´ A. Orlová, K. Kuchaˇrová, A. Dlouhy, Mater. Sci. Eng. A483–484 (2008) 109– 112. J. Lapin, Intermetallics 14 (2006) 115–122. L.M. Hsiung, T.G. Nieh, B.W. Choi, J. Wadsworth, Mater. Sci. Eng. A329–331 (2002) 637–643. F. Povolo, J. Mater. Sci. 20 (1985) 2005–2010.