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Mechanical behavior of submicron-grained γ-TiAl-based alloys at elevated temperatures
Intermetallics 9 (2001) 559–569 www.elsevier.com/locate/intermet
Mechanical behavior of submicron-grained g-TiAl-based alloys at elevated temperatures R. Bohna,*, T. Klassena, R. Bormanna,b a Institute for Materials Research, GKSS Research Center Geesthacht, D-21502 Geesthacht, Germany Department of Material Science and Technology, Technical University Hamburg-Harburg, D-21073 Hamburg, Germany
b
Received 2 March 2001; received in revised form 6 April 2001; accepted 24 April 2001
Abstract Submicron-grained intermetallic compounds based on g-TiAl were prepared by high-energy milling and hot isostatic pressing. At temperatures above 500 C, the flow stress is strongly reduced and work-hardening completely disappears. Compression tests performed at temperatures between 500 and 900 C reveal a marked strain rate sensitivity of the flow stress, suggesting superplasticity to occur. This could be confirmed by tensile straining of Ti–45Al–2.4Si samples, resulting in elongations of up to 175% at 800 C. Small silicide particles (d100 nm) of the type Ti5(Si,Al)3, embedded in the grain boundaries of the g-TiAl matrix, impede a coar: sening of the microstructure. However, at strain rates above " ¼ 103 s1 , these dispersoids are suggested to promote the formation of voids and to reduce the overall deformability. At 800 C, an apparent activation energy of Qapp ¼ 351 kJ/mol can be derived. Superplastic behavior at 800 C is accomplished by grain boundary sliding accommodated by diffusional processes inside the g-TiAl phase. Thus, the high temperature deformation mode is similar to the mechanisms found for more conventionally grained TiAl alloys at deformation temperatures 51000 C. # 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Nanostructured intermetallics; A. Titanium aluminides, based on TiAl; C. Superplastic forming; C. Mechanical alloying and milling
1. Introduction Many studies on nano- and submicron-grained materials have dealt with the mechanical properties at room temperature, in order to explore the effect of exceptional microstructural refinement on strength and ductility (e.g. [1–4]). However, little is known about the mechanical behavior at elevated temperatures 50.5 Tm , with Tm denoting the absolute melting temperature. One reason may be the low thermal stability of ultrafinegrained materials, giving rise to rapid grain growth upon increasing the temperature. Consequently, the few existing studies are mostly concerned with multiphase metallic systems [5–8] or ceramics [9], that are less prone to grain coarsening. Furthermore, specimens with flawfree microstructures are required, for it was shown that pores — which are a common feature of many powdermetallurgically processed samples — influence creep properties via increasing the diffusion rates [10]. Creep tests performed on nanocrystalline Ni–P and Fe–B–Si * Corresponding author.
samples result in small stress exponents and low activation energies for creep [6,7], suggesting deformation mechanisms that are based on grain-boundary-diffusion assisted processes. However, quantitative analysis of the creep data obtained on nanocrystalline Cu and Pd [8] demonstrates, that at least for moderate temperatures up to 0.33 Tm , the resulting strain rates are much lower than those according to the equation for Coble creep [11]. There are also some contradictory results concerning the grain size dependence of the deformation strain rate. Whereas Kong et al. show that nanocrystalline samples deform much faster than their submicron-grained counterparts [5,7], Sanders et al. cannot observe any difference between the creep rates of ultrafine- or coarse-grained Cu and Pd samples [8]. In the latter case, a large fraction of low energy interfaces like twin- and low angle boundaries was supposed to be responsible for the poor diffusivity and the lack of grain boundary sliding. Generally, ultrafine-grained microstructures with random high angle boundaries should promote deformation at elevated temperatures and lead to easy workable, low strength materials. This was recently demonstrated by
0966-9795/01/$ - see front matter # 2001 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(01)00039-5
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hot working of submicron-grained g-TiAl based compounds [12,13]. Isothermal forging, extrusion as well as superplastic deformation could be performed at comparatively low temperatures or stresses. The purpose of this study is to elucidate the underlying deformation mechanisms. Ameyama et al. have shown [14], that the superplastic behavior of a Ti–48Al (in at.%) compound with a mean grain size of about 1 mm is based on grain boundary sliding accommodated by lattice diffusion inside the g-TiAl phase. That kind of mechanisms does not differ from the deformation mode described for coarser grained material, e.g. by Mishra et al. [15]. In this study, much finer microstructures are examined, and the question arises, whether the close meshed net of grain boundaries contributes to deviations in the deformation behavior. In order to increase the resistance against grain growth at elevated temperatures, the gTiAl matrix of the compounds investigated is reinforced with a fine dispersion of silicide particles.
2. Experimental Prealloyed gas-atomized powders of composition Ti– 48.9Al and Ti–37.5Si as well as pure silicon powders were mixed in different proportions and high-energy milled in an 8 l attrition mill of Zoz GmbH. Consolidation of the powders to fully dense compacts was accomplished by hot isostatic pressing (HIP) at temperatures between 800 and 1000 C. To minimize any contamination of the powders with atmospheric impurities, milling, transport and storage of the powders were conducted under vacuum or controlled gas atmospheres (high purity Ar (99.9999%) or Ar–5 vol.% H2). More detailed information about the powder processing may be taken from Refs. [16,17]. Microstructural investigations were performed using a Philips CM 200 transmission electron microscope combined with a Link ISIS energy dispersive X-ray analyser. The average grain size of the consolidated samples was determined from the micrographs by applying the linear intercept method. The room temperature microhardness was measured on metallographically polished samples under a load of 1 N. Data points represent the average of 40 measurements each. Compression and tensile tests were carried out on spark-eroded and ground specimens. For compression testing, cylindrical specimens with a diameter of 4 mm and a length of 8 mm were used. The Zwick 1478 universal testing machine was equipped with a three-zone furnace to ensure a constant and uniform heating of both the specimen and the pistons. In order to minimize the influence of oxidation, the compression tests were run in an atmosphere of flowing Ar with a purity of 99.998%. The compressive strain was recorded via 3 alumina rods, directly pressed against the specimen or the upper die, respectively, and connected to a LVDT outside the heating zone of the furnace. The maximum
deformation detectable was limited to 1 mm, corresponding to a strain of 12.5%. Whereas compression tests were run with constant velocity, tensile tests on a Zwick 1484 testing machine were performed using constant strain rates. The corresponding specimens had a gauge length of 22 mm. Tensile tests could be carried out only in air.
3. Results 3.1. Microstructure The intermetallic/ceramic compounds, produced by high-energy-milling and HIP, are very homogeneous and consist of equi-axed g-TiAl-grains and x-Ti5(Si,Al)3particles. Depending on the adjusted volume fraction of the silicide phase, two types of microstructures may be distinguished. (a) Intermetallic/ceramic composites with a moderate silicide content of about 9.6 vol.% (as in Ti–45Al–2.4Si) or 13.1 vol.% (as in Ti–46Al–5Si) reveal a bimodal grain size distribution, i.e. submicron grains of the majority phase g-TiAl are surrounded by clearly smaller particles of the x-Ti5(Si,Al)3-phase (Fig. 1a). Different HIP-temperatures do not influence this overall appearance. Volume fractions and arrangement of the phases involved do not alter, only the phase specific mean grain sizes vary, but without changing their mutual ratio. An estimation of the finest grain size achievable under the applied HIP conditions is given in Ref. [18]. (b) Highly silicide containing compounds like Ti– 36Al–10Si, with about 30 vol.% of x-Ti5(Si,Al)3, do not show any difference between the mean grain sizes of the g-TiAl phase and the silicide phase (Fig. 1b). The silicide content is however still low enough to prevent the formation of a continuous ceramic network, i.e. the Ti5(Si,Al)3-grains may be regarded as isolated inclusions between the g-TiAl crystallites. Silicon-free binary TiAl-compounds reveal a globular microstructure of the g phase with small 2 grains precipitated at the triple points of g grain boundaries [19]. These alloys were included in the study as a submicrongrained, but silicon-free reference material. The microstructural data of all samples investigated were recently published in [4,16] along with their impurity contents. Considering the multistep powdermetallurgical processing of the samples, it should be noted that the amount of contamination by interstitial atoms like O (0.3 at.%) and N (0.1 at.%) is very low. Nevertheless, the material does not reach the purity of ingots. As the content of O and N probably exceeds the solubility limit of the g-TiAl phase, small amounts of oxides and nitrides should be present within the microstructure. A more detailed description of the microstructure is included in Refs. [16,17].
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Fig. 1. TEM micrographs showing the microstructure of TiAlSi compounds in the as-HIP’ed condition; (a) Ti–45Al–2.4Si, consolidated at 875 C; (b) Ti–36Al–10Si, compacted at 1000 C. Further details are provided in Ref. [16].
3.2. Temperature dependence of yield stress At low temperatures, submicron-grained alloys based on g-TiAl are characterized by high values of hardness and compressive yield strength [20]. As previously reported [4], this behavior can be explained by the clasp sical Hall–Petch relationship, resulting in a 1/ ddependence of grain size d on yield strength y , i.e. dislocation glide and mechanical twinning as the prevailing deformation mechanisms are hampered by the finemeshed net of grain-boundaries present in submicrongrained alloys. Consequently, further grain refinement leads to a continued increase of hardness and strength. In the high temperature range above 500 C, these relations change completely, i.e. ultrafine-grained materials become very soft. As demonstrated in Fig. 2, there is a
Fig. 2. Flow stress after 1.25% plastic compression strain, 1.25, in dependence of the deformation temperature for silicon-free alloys; ML, AM and CF refer to Mechanically Alloyed [19], high energy Attrition Milled (this work) and Cast/Forged material [21].
grain size-dependent inversion of yield strength, which means that the smaller the grain size of the specimen, the stronger and sharper the drop in yield strength upon raising the temperature. Silicon-doped g-TiAl-based compounds with a fraction of up to 32 vol.% Ti5(Si,Al)3 show a similar behavior as the binary TiAl alloys (Fig. 3). Grain refinement causes a reduction of flow stress, suggesting favorable conditions for hot-working of ultrafine-grained TiAlSi compounds. 3.3. Strain hardening Increasing the test temperature leads to a reduction of strain hardening. As demonstrated in Fig. 4, a binary TiAl sample with a mean grain size of 670 nm deforms already at 700 C under constant stress, i.e. the dynamic sample loading with a given velocity of the crosshead is comparable to a static loading under a corresponding constant stress. Thus, the dynamic compression tests are
Fig. 3. Flow stress 1.25 vs. temperature for Ti–46Al–5Si and Ti– 36Al–10Si with a silicide fraction of 13 or 32 vol.%, respectively.
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Fig. 4. Compression curves obtained on specimens of composition Ti– 48.9Al with a mean grain size of 670 nm.
formally quite similar to the steady state regime of creep tests.1 Specimens with smaller grain sizes show a lack of strain hardening at still lower temperatures. 3.4. Strain rate dependence of flow stress Submicron-grained TiAlSi compounds are characterized by a marked strain rate sensitivity at temperatures 5500 C. In Fig. 5, the flow stress of differently finegrained specimens of composition Ti–45Al–2.4Si is depicted as a function of the applied compression rate at a temperature of 800 C. The coefficient of strain rate sensitivity, m, that describes the slope of the respective curves, varies between 0.3 and 0.55. Values of that magnitude are indicative of superplastic deformation behavior. However, in order to actually demonstrate the potential of submicron-grained g-TiAl-based alloys for superplastic deformation at comparatively low temperatures, tensile tests are required. The corresponding results are presented in Section 3.6. The strong strain rate dependence of submicrongrained TiAlSi compounds has to be kept in mind if the temperature dependence of the flow stress (Figs. 2 and 3) is interpreted. Especially, from a high flow stress (under the strain rate conditions applied), it does not automatically follow that a good creep resistance will be observed. At constant stress, decreasing the grain size allows for higher strain rates to be adjusted in order to reach the maximum strain rate sensitivity. In the present case, the bisection of the mean grain size admits about a fivefold increase of the strain rate.
1 The expression ‘‘creep test’’ may be somewhat irritating, particularly since deformation rates in the range of 104–103 s1 are considered. However, it should be pointed out that similar deformation curves as in Fig. 4 are expected by applying a constant stress instead of a constant crosshead velocity.
Fig. 5. Strain rate dependence of flow stress 3.0 (stress after 3.0% true plastic compression strain) at 800 C for compounds of Ti–45Al– 2.4Si with different mean grain sizes. Fig. 4 represents a completion of Fig. 5 in Ref. [13].
Similar diagrams as in Fig. 5 are attained if the grain size is kept constant, but the test temperature is changed. From Fig. 4 in Ref. [13], it is obvious that grain refinement is comparable to a temperature increase with respect to the deformation behavior. This fact is of great technological interest, as it allows hot-working of submicron-grained TiAl alloys at reduced temperatures. 3.5. Microstructural stability Nano- and submicron-grained microstructures are prone to grain growth due to the large amount of stored grain boundary enthalpy. Therefore, it has to be investigated to what extent the mechanical data presented in the previous sections still refer to the as-HIP’ed microstructures shown in Fig. 1. The grain size of the test specimens was controlled twice. Firstly, some samples were annealed at the deformation temperatures of interest and subsequently cooled down to room temperature. The dwell time chosen for these heat treatments was 5 h, which is about 2 h more than the preheating time necessary to adjust a steady temperature profile within the ceramic pistons of the testing machine. This gives a reference for the microstructure at the start of the deformation procedure. Secondly, the deformed specimens were investigated in order to get informations about microstructural changes occurring upon deformation. As mentioned before, room temperature hardness and yield strength of the material are closely related to the grain size via the Hall-Petch correlation. Therefore, instead of the time-consuming preparation of TEM foils, it was sufficient to perform hardness measurements in order to get a first impression of microstructural changes. Occasional spot checks carried out in the TEM revealed that these microstructural changes were in fact confined to grain growth processes. The
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Fig. 7. Room temperature hardness, obtained on specimens of composition Ti–45Al–2.4Si after 12.5% compressive deformation at 800 C, in dependence of the applied strain rate. Hardness data referring to tensile strained specimens ‘‘B’’ and ‘‘F’’ is also included (see Section 3.6 for more details).
Fig. 6. (a) Room temperature hardness as a function of annealing temperature for three g-TiAl based compounds with different fractions of the silicide phase Ti5(Si,Al)3, according to 0!10!13 vol.%. The resulting grain growth at temperatures corresponding to the applied HIP-temperatures is calculated on basis of the Hall–Petch parameters given in Ref. [4]. (b) Temperature-dependent flow stress 1.25 of the same compounds.
results of investigations on the microstructural stability are shown in Figs. 6 and 7. In Fig. 6a, the hardness at room temperature is plotted against the chosen annealing temperature. As long as a high hardness level is retained, grain growth does not occur. The applied HIPtemperatures are also indicated. As expected, the stability of the microstructure at the respective HIP-temperature is enhanced by increasing the amount of the silicide phase (0!10!13 vol.%). In Fig. 6b, the yield stresses of the different compounds are shown as a function of the deformation temperature. The three curves widely overlap. In combination, Fig. 6a and b show that the marked drop of yield strength occurs already before the onset of considerable grain growth, at least in the case of the silicon-doped material. Accordingly, the decrease of flow stress is rather caused by a change of the deformation mode than by microstructural instabilities of the test specimens. This is an important result with respect to technological applications: The low flow stress due to the fine grain size and the silicide-dispersion-induced stability of the microstructure admit favorable hot-working conditions for the shaping of components.
After compression straining at 800 C, specimens of alloy type Ti–45Al–2.4Si do not show any difference in hardness between the unstrained or the deformed state, independent of the strain rate applied (Fig. 7). This allows the conclusion, that the initial microstructure and the respective mechanical properties did not significantly change. However, it also has to be kept in mind that the total compressive strain amounts to only 12.5%, which might be not enough for the start of strain-induced reactions like recrystallization. Regarding tensile tests, the preconditions are somewhat different, as discussed in Section 4. 3.6. Tensile straining Tensile tests at 800 C were performed on the basis of the mechanical data derived from compression straining of Ti–45Al–2.4Si. The experiments were launched to validate the supposed potential for superplastic deformability and to verify the results obtained by compressive loading. The results are summarized in Table 1, and may be described as follows: : 1. At strain rates "43:2103 s1 elongations "5175% were reached (Fig. 8). Larger deformations could not be managed due to the limited length of the : furnace. By increasing the strain rate to "46:4 103 s1 the specimens fail after elongations of 103 or 113%, respectively. 2. The coefficients of the strain rate sensitivity are comparable to those obtained by compression testing. Increasing the strain rate leads to higher strain rate sensitivities. : 3. Simultaneously to an increase of " and m, the fracture strain drops. However, necking of the samples is not observed.
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Table 1 : Coefficient of strain rate sensitivity, fracture strain, porosity and resulting grain size of Ti–45Al–2.4Si specimens deformed at different strain rates. "1 : and "2 are the strain rates applied during the strain rate cycle tests
Fig. 9. TEM micrograph, taken from the neck region of specimen ‘‘B’’ after a deformation of 175%. While the grain shape did not change, slight grain coarsening from 170!248 nm can be observed.
Fig. 8. Tensile samples of the compound Ti–45Al–2.4 Si after deformation at 800 C in air. An undeformed specimen is shown for comparison.
4. The microstructure has conserved its equi-axed grain character after deformation (Fig. 9). Contrary to compression loaded specimens, limited grain growth occurs. However, within the strain rate range investigated, increasing the deformation rate leads to a reduction of grain growth. 5. Measurements of the room temperature hardness were carried out in both the neck and the head region of deformed specimens. The results are presented in Table 1 and Fig. 7. Whereas the hardness of the specimen heads did not show any significant changes in comparison to
the as-HIP’ed state, the neck regions have become considerably softer, irrespective of the strain rate applied. As mentioned above, grain growth proved to depend on the strain rate and thus can not explain the loss of hardness. Instead, the softening is attributed to the formation of voids, especially in fast strained specimens (Fig. 10). The accumulation of pores near surface notches indicates that local stress concentrations largely contribute to cavitation.
4. Discussion The high temperature mechanical behavior of submicron-grained g-TiAl-based compounds is essentially characterized by the strongly reduced flow stress. This property already allowed for rather low consolidation temperatures during powder processing and implicitly
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napp
Fig. 10. Evolution of voids in specimen ‘‘F’’ upon tensile straining at 800 C. The SEM micrograph is taken at the edge of the central neck region of tensile specimen ‘‘F’’. The surface roughness leads to a nonuniform distribution of voids. A lot of pores have accumulated at the tip of a notch (arrowed), whereas other regions are essentially free of voids (encircled).
contributed to the fine-grained microstructure. In order to elucidate the mechanisms of high temperature deformation, the stress-strain curves will be further analyzed. As indicated above, the compression tests can be treated like static deformation tests, i.e. the compressive : strain rate " — which is actually a given parameter — was analysed in dependence of the plateau stress 3:0 arising after an initial plastic deformation of three per: cent. Generally, " is a function of temperature T, stress and quite a lot of microstructural parameters Si , like grain size, grain shape, volume fractions and distribution of phases involved etc.: : " ¼ fðT; ; Si Þ
ð1Þ
In the simplest case, these parameters do not interact. Assuming that essentially one process is dominating, the analysis may be based on the equation n Q : " ¼ A
exp E RT
ð2Þ
E denotes the Young modulus, n is the stress exponent and Q refers to the activation energy of the deformation step that is rate-controlling. The influence of the microstructural parameters is represented by the prefactor A. The slope of the curves shown in Fig. 5 leads to the respective coefficient of the strain rate sensitivity, m, which is reciprocal to the apparent stress exponent
: dln" 1 ¼ ¼ dln T¼const: m
565
ð3Þ
For small stresses, napp is about equal to n as long as the prefactor A does not depend on stress. At larger stresses, napp exceeds n. In advance to the analysis, several aspects concerning the microstructural stability upon high temperature deformation have to be assessed. 1. The mean grain size before and after deformation was evaluated by hardness measurements. However, this procedure only offers a selective shot of the truth. For example, it is possible that a nonvariant grain size is the result of both grain growth and dynamic recrystallization, that balance each other. This assumption is supported by studies of Imayev et al. [22] and Koeppe et al. [23], suggesting that grain refinement or grain growth prevails depending on whether a certain strain rate or stress level is exceeded or not. The results of the tensile tests performed within this study tend to support this view. For the Ti–45Al–2.4Si compound with a mean grain size of 170 nm, this neutral strain rate is obviously larger than 3.2103 s1 at 800 C. Thus, it may be explained why the slowly strained specimen ‘‘B’’ shows more intense grain coarsening than specimen ‘‘F’’. In the case of the compression tests, it may be argued that the strain level attained upon deformation should be too low to observe similar effects. 2. In strain rate cycle tests performed both in tension and under compression, the doubling of the strain rate does not lead to a spontaneous change of flow stress. Instead, the corresponding stress level is reached rather gradually, which is indicative of microstructural changes occurring inside the specimen. Apart from variations of the grain size, the number and arrangement of dislocations, twins etc. could account for this behavior. Consequently, the structure-sensitive prefactor A in Eq. (2) should not virtually be treated as a constant, and napp should deviate from n. However, the actual differ: ence must be very small, as strain rate changes from " i : to " iþ1 (and back) lead to the same plateau stresses that are reached if the deformation tests are performed with : a strain rate "iþ1 right from the beginning. Summarizing, the question of microstructural stability can be regarded as rather uncritical with respect to the performed analysis. In this study, napp amounts to values between 1.8 and 3.3. Superplasticity based on grain boundary sliding requires n=2 [24]. The unchanged grain shape after a tensile strain of 175% provides strong evidence for grain boundary sliding to play an important role. In contrast to this, deformation modes based on diffusional flow like Coble [11] or Nabarro– Herring creep [25] are characterized by n ¼ 1 and would lead to an elongation of the crystallites. As this is not observed, it is concluded that napp does not significantly overestimate n.
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Grain boundary sliding as an essential part of the superplastic deformation must always be supported by accommodation processes that prevent the formation of voids at grain triple junctions. The temperature dependence of deformation can provide further information about the nature of the assisting mechanism. By plotting : " versus 1/T, the apparent activation energy may be derived according to 0 1 : B dln" C Qapp ¼ R @ ð4Þ A 1 dln T ¼const:
small particles require higher stresses for dislocation generation and multiplication. More information about the deformation mode can be drawn from the grain size dependence of the strain rate. The microstructure-sensitive parameter A may be expressed in the form A / dp
ð5Þ
In Fig. 11, the modulus-compensated strain rate under compression is plotted against 1/T. The slope of the curves leads to a value of Qapp ¼ 351 kJ/mol, irrespective of whether the compounds Ti–45Al–2.4Si or Ti– 36Al–10Si are regarded. Obviously, the same mechanisms are operative in both materials. This is not surprising, assuming that grain boundary sliding is assisted by processes controlled by only the g-TiAl phase. Though Frommeyer et al. specify an apparent activation energy of 350 kJ/mol also for the deformation of monolythic Ti5Si3 at temperatures between 1000 and 1300 C [27], it seems unlikely that the hard x-Ti5(Si,Al)3 phase is codeformed within a composite containing the softer gTiAl phase. This assumption is substantiated by studies of Wang et al. on powder-metallurgically processed specimens of composition Ti–45Al–2.7Si, showing that the silicide particles are free of dislocations after deformation at 700 C [28]. The same should be valid with respect to the still finer silicide grains of this study, especially as
The grain size exponent p results from the slope of a : ln" vs. lnd plot (Fig. 12). At T ¼ 800 C and ¼ 150 MPa, a grain size exponent of p ¼ 2:2 is obtained. Classifying the experimentally derived parameters, Qapp , napp and p the deformation mechanism may be characterized as follows: The apparent activation energy of 351 kJ/mol is somewhat higher than the reported activation energy of 250–290 kJ/mol for the titanium self diffusion in g-TiAl [29,30]. Nevertheless, it compares well to the published activation energies of comparatively coarse-grained g-TiAl alloys (Table 2), that were deformed at similar or even higher (>1000 C) temperatures. In connection with a grain size parameter closer to 2, it is concluded that processes based rather on volume diffusion than on grain boundary diffusion support the grain boundary sliding of submicrocrystalline TiAlSi compounds. A corresponding model is sketched in Fig. 13. The diffusional flow inside the g-TiAl grains may be provided by dislocation climb, for example. The stress and grain size exponents are very similar to the values published by Ameyama et al. [14] and Mishra et al. [15] for the superplastic deformation of gTiAl-based alloys with grain sizes of about 1 mm [14] or 20 mm [15], respectively. Thus, the reduction of the
Fig. 11. True compressive strain rate, compensated for the temperature dependence [26] of the Young’s modulus, as a function of the reciprocal temperature for compounds of a type Ti–45Al–2.4Si and Ti–36Al–10Si. The analysis was performed at a stress level of 150 MPa.
Fig. 12. True compressive strain rate in dependence of the mean grain size, demonstrated for samples of composition Ti–45Al–2.4Si. The grain size exponent of the deformation was derived for a temperature of 800 C and a stress level of 150 MPa.
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Table 2 Stress exponents and apparent activation energies obtained from high temperature deformation experiments performed on g-TiAl based alloys [14–15], [19], [31–37]. a and b refer to tensile or compression tests, respectively. The alloys marked by * could be superplastically deformed under the conditions specified
mean grain size down to the low submicron range does not cause a change in the prevailing deformation mechanism, it rather decreases the lower boundary of the temperature regime where these mechanisms work. The results of this study are in some contrast to recently published work of Imayev et al. [36]. According to a specified activation energy of 180–200 kJ/mol for the superplastic deformation of submicron-grained TiAl-based alloys at moderate temperatures of about 750–850 C, the authors claim deformation to be accomplished by grain boundary sliding assisted by dif-
fusion along the grain boundaries. The reason for this discrepancy cannot be ascertained. However, it should be noted, that in a corresponding temperature range, similar specimens as those investigated in [36] require an activation energy of 335 kJ/mol for grain growth [39], which rather favours volume instead of grain-boundary diffusion to play a dominant role. The formation of voids upon superplastic deformation of g-TiAl alloys is a frequently observed phenomenon [40–42]. Reducing the grain size was shown to be an appropriate countermeasure [41]. In the present case, a
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Fig. 13. Model for the superplastic deformation of submicron-grained Ti–45Al–2.4Si: grain-boundary sliding, accommodated by dislocation creep inside the g-TiAl crystals (picture taken from Refs. [38]).
lot of pores are visible in the vicinity of surface notches. Thus it is concluded that stress-induced cavitation plays an important role. On a submicron scale, hard silicide particles are supposed to be responsible for the occurrence of local stress concentrations leading to the formation of voids all over the neck of the sample. This may be illustrated by replacing grain 3 of Fig. 13 by a silicide particle impermeable for dislocations. In general, the grain triple points can be identified as preferential nucleation sites for the formation of cracks [43]. According to Stroh [44], cracks may be initiated if the shear stress exceeds a critical value of c ¼
12G L
ð6Þ
with denoting the surface energy and L representing the length of the glide area. From Eq. (6) it is obvious that grain refinement allows for higher critical stresses. On the other hand, a lot of triple points are occupied by silicide particles that could reduce the surface energy and thus balance out the positive influence of a small grain size. In this respect, an even finer distribution of smaller silicides might be useful.
5. Summary and conclusions 1. Multiphase g-TiAl-based compounds with a very fine-grained microstructure were produced by high energy milling and hot isostatic pressing. The grain size of the g-TiAl matrix amounts to about 160–480 nm,
whereas the size of precipitated x-Ti5(Si,Al)3 particles varies between 80 and 190, depending on the HIP conditions. 2. At temperatures above 500 C, these alloys become very soft. The smaller the grain size, the stronger the drop of flow stress. Deformation occurs without any strain hardening. 3. During deformation at high temperatures, the finely dispersed silicide particles prevent grain growth as long as the HIP temperature is not exceeded. 4. The flow stress is rather sensitive to the applied strain rate. Superplasticity becomes feasible at 800 C, allowing elongations 5175%. The apparent activation energy of deformation was assessed to be Qapp ¼ 351 kJ/ mol. In connection with a stress exponent around n ¼ 2 and a grain size exponent of p ¼ 2:2, it is concluded that the superplastic deformation is accomplished by grain boundary sliding accommodated by diffusional processes inside the g-TiAl grains. Such a mechanism does not differ from the deformation modes described for conventionally grained materials deformed at temperatures 51000 C. 5. The silicide phase is not directly involved in the process of deformation. In particular, Ti–36Al–10Si compounds with about 32 vol.% of the silicide phase reveal the same activation energy as alloys of composition Ti–45Al–2.4 Si with only 9.6 vol.% Ti5(Si,Al)3. In both cases only the g-TiAl phase seems to contribute to the overall deformation. However, changes are expected if the silicide content is high enough to form a continuous network within the g-phase. 6. Increasing the tensile strain rate supports the formation of stress induced voids that reduce the overall deformability. The silicide particles are supposed to promote cavitation.
Acknowledgements This work has been supported by the Deutsche Forschungsgemeinschaft (German Science Foundation) within the scope of Sonderforschungsbereich 371. Thanks are due to Dr. R. Gerling for providing the powders for high energy milling. The experimental help of U. Lorenz and R. Behn is gratefully acknowledged.
References [1] Gertsman VY, Hoffmann M, Gleiter H, Birringer R. Acta Metall Mater 1994;42:3539. [2] Sanders PG, Eastman JA, Weertman JR. Acta Mater 1997; 45:4019. [3] Koch CC, Morris DG, Lu K, Inoue A. MRS Bulletin 1999;24:54. [4] Bohn R, Klassen T, Bormann R. Acta Mater 2001;49:299. [5] Wang DL, Kong QP, Shui JP. Scripta Metall Mater 1994;3:47. [6] Deng J, Wang DL, Kong QP, Shui JP. Scripta Metall Mater 1995;32:349.
R. Bohn et al. / Intermetallics 9 (2001) 559–569 [7] Xiao M, Kong QP. Scripta Metall Mater 1997;36:299. [8] Sanders PG, Rittner M, Kiedaisch E, Weertman JR, Kung H, Lu YC. Nanostr Mater 1997;9:433. [9] Hahn H, Averback RS. Nanostr Mater 1992;1:95. [10] McMeeking RM, Kuhn LT. Acta Mater 1992;40:961. [11] Coble RL. J Appl Phys 1963;34:1679. [12] Bohn R, Fanta G, Bormann R. In: Morris DG et al., editors. Intermetallics and superalloys. Weinheim, Germany: WileyVCH, 2000. p. 271. [13] Fanta G, Bohn R, Dahms M, Klassen T, Bormann R. Intermetallics 2001;9:45. [14] Ameyama K, Uno H, Tokizane M. Intermetallics 1994;2:315. [15] Mishra RS, Lee WB, Mukherjee AK, Kim Y-W. In: Kim Y-W et al., editors. Gamma titanium aluminides. Warrendale, USA: TMS, 1995. p. 571. [16] Bohn R, Fanta G, Klassen T, Bormann R. J Mater Res, in press. [17] Bohn R, PhD thesis, Technische Universita¨t Hamburg-Harburg, 1999 (in German). [18] Bohn R, Klassen T, Bormann R. Adv Eng Mater 2001;3:238. [19] Oehring M, Appel F, Pfullmann Th, Bormann R. Appl Phys Lett 1995;66:941. [20] Bohn R, Oehring M, Pfullmann T, Appel F, Bormann R. In: Suryanarayana C et al., editors. Processing and Properties of Nanocrystalline Materials. Warrendale, USA: TMS, 1996. p. 355. [21] Appel F, Lorenz U, Sparka U, Wagner R. In: Oikawa et al., editors. Strength of metals — Proc. of the 10th Int. Conference ICSMA 10. Sendai, Japan: JIM, 1994. p. 341. [22] Imayev RM, Imayev VM, Salishchev GA. J Mater Sci 1992; 27:4465. [23] Koeppe C, Bartels A, Clemens H, Schretter P, Glatz W. Mater Sci Eng 1995;A201:182. [24] Lu¨thy H, White RA, Sherby OD. Mater Sci Eng 1979;39:211. [25] Herring C. J Appl Phys 1950;21:437. [26] Schafrik RE. Metall Trans 1977;8A:1003.
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[27] Frommeyer G, Rosenkranz R, Lu¨decke CZ. Metallkd 1990; 8:307. [28] Wang G-X, Dogan B, Hsu F-Y, Klaar H-J, Dahms M. Metall Trans 1995;26A:691. [29] Mishin Y, Herzig C. Acta Mater 2000;48:589. [30] Kroll S, Mehrer H, Stolwijk N, Herzig C, Rosenkranz R, Frommeyer GZ. Metallkd 1992;83:591. [31] Seetharaman V, Lombard CM. In: Kim Y-W et al., editors. Microstructure/property relationships in titanium aluminides and alloys. Warrendale, USA: TMS, 1991. p. 237. [32] Martin PL, Mendiratta MG, Lipsitt HA. Metall Trans A 1983; 14A:2170. [33] Kampe SL, Bryant JD, Christodoulou L. Metall Trans A 1991; 22A:447. [34] Cheng SC, Wolfenstine J, Sherby OD. Metall Trans A 1992; 23A:1509. [35] Wang JG, Nieh TG. Intermetallics 2000;8:737. [36] Imayev RM, Salishchev GA, Senkov ON, Imayev VM, Shagiev MR, Gabdullin NK, Kuznetsov AV, Froes FH. Mater Sci Eng A 2001;A300:263. [37] Rommerskirchen M. PhD thesis, RWTH Aachen, 1997 (in German). [38] Sherby OD, Wadsworth J. Progr Mater Sci 1989;33:169. [39] Senkov ON, O¨vec¸oglu ML, Srisukhumbowornchai N, Froes FH. Nanostr Mater 1998;10:935. [40] Lee WB, Yang HS, Kim Y-W, Mukherjee AK. Scripta Metall Mater 1993;29:1403. [41] Imayev VM, Imayev RM, Salishchev GA. Mater Sci Eng 1996; A208:226. [42] Appel F, Clemens H, Glatz W, Wagner R. MRS Symp Proc 1997;460:195. [43] Edington JW, Melton KN, Cutler CP. Prog Mater Sci 1976; 21(2):61. [44] Stroh AN. Proc Roy Soc 1934;A223:404.
Mechanical behavior of submicron-grained g-TiAl-based alloys at elevated temperatures R. Bohna,*, T. Klassena, R. Bormanna,b a Institute for Materials Research, GKSS Research Center Geesthacht, D-21502 Geesthacht, Germany Department of Material Science and Technology, Technical University Hamburg-Harburg, D-21073 Hamburg, Germany
b
Received 2 March 2001; received in revised form 6 April 2001; accepted 24 April 2001
Abstract Submicron-grained intermetallic compounds based on g-TiAl were prepared by high-energy milling and hot isostatic pressing. At temperatures above 500 C, the flow stress is strongly reduced and work-hardening completely disappears. Compression tests performed at temperatures between 500 and 900 C reveal a marked strain rate sensitivity of the flow stress, suggesting superplasticity to occur. This could be confirmed by tensile straining of Ti–45Al–2.4Si samples, resulting in elongations of up to 175% at 800 C. Small silicide particles (d100 nm) of the type Ti5(Si,Al)3, embedded in the grain boundaries of the g-TiAl matrix, impede a coar: sening of the microstructure. However, at strain rates above " ¼ 103 s1 , these dispersoids are suggested to promote the formation of voids and to reduce the overall deformability. At 800 C, an apparent activation energy of Qapp ¼ 351 kJ/mol can be derived. Superplastic behavior at 800 C is accomplished by grain boundary sliding accommodated by diffusional processes inside the g-TiAl phase. Thus, the high temperature deformation mode is similar to the mechanisms found for more conventionally grained TiAl alloys at deformation temperatures 51000 C. # 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Nanostructured intermetallics; A. Titanium aluminides, based on TiAl; C. Superplastic forming; C. Mechanical alloying and milling
1. Introduction Many studies on nano- and submicron-grained materials have dealt with the mechanical properties at room temperature, in order to explore the effect of exceptional microstructural refinement on strength and ductility (e.g. [1–4]). However, little is known about the mechanical behavior at elevated temperatures 50.5 Tm , with Tm denoting the absolute melting temperature. One reason may be the low thermal stability of ultrafinegrained materials, giving rise to rapid grain growth upon increasing the temperature. Consequently, the few existing studies are mostly concerned with multiphase metallic systems [5–8] or ceramics [9], that are less prone to grain coarsening. Furthermore, specimens with flawfree microstructures are required, for it was shown that pores — which are a common feature of many powdermetallurgically processed samples — influence creep properties via increasing the diffusion rates [10]. Creep tests performed on nanocrystalline Ni–P and Fe–B–Si * Corresponding author.
samples result in small stress exponents and low activation energies for creep [6,7], suggesting deformation mechanisms that are based on grain-boundary-diffusion assisted processes. However, quantitative analysis of the creep data obtained on nanocrystalline Cu and Pd [8] demonstrates, that at least for moderate temperatures up to 0.33 Tm , the resulting strain rates are much lower than those according to the equation for Coble creep [11]. There are also some contradictory results concerning the grain size dependence of the deformation strain rate. Whereas Kong et al. show that nanocrystalline samples deform much faster than their submicron-grained counterparts [5,7], Sanders et al. cannot observe any difference between the creep rates of ultrafine- or coarse-grained Cu and Pd samples [8]. In the latter case, a large fraction of low energy interfaces like twin- and low angle boundaries was supposed to be responsible for the poor diffusivity and the lack of grain boundary sliding. Generally, ultrafine-grained microstructures with random high angle boundaries should promote deformation at elevated temperatures and lead to easy workable, low strength materials. This was recently demonstrated by
0966-9795/01/$ - see front matter # 2001 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(01)00039-5
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hot working of submicron-grained g-TiAl based compounds [12,13]. Isothermal forging, extrusion as well as superplastic deformation could be performed at comparatively low temperatures or stresses. The purpose of this study is to elucidate the underlying deformation mechanisms. Ameyama et al. have shown [14], that the superplastic behavior of a Ti–48Al (in at.%) compound with a mean grain size of about 1 mm is based on grain boundary sliding accommodated by lattice diffusion inside the g-TiAl phase. That kind of mechanisms does not differ from the deformation mode described for coarser grained material, e.g. by Mishra et al. [15]. In this study, much finer microstructures are examined, and the question arises, whether the close meshed net of grain boundaries contributes to deviations in the deformation behavior. In order to increase the resistance against grain growth at elevated temperatures, the gTiAl matrix of the compounds investigated is reinforced with a fine dispersion of silicide particles.
2. Experimental Prealloyed gas-atomized powders of composition Ti– 48.9Al and Ti–37.5Si as well as pure silicon powders were mixed in different proportions and high-energy milled in an 8 l attrition mill of Zoz GmbH. Consolidation of the powders to fully dense compacts was accomplished by hot isostatic pressing (HIP) at temperatures between 800 and 1000 C. To minimize any contamination of the powders with atmospheric impurities, milling, transport and storage of the powders were conducted under vacuum or controlled gas atmospheres (high purity Ar (99.9999%) or Ar–5 vol.% H2). More detailed information about the powder processing may be taken from Refs. [16,17]. Microstructural investigations were performed using a Philips CM 200 transmission electron microscope combined with a Link ISIS energy dispersive X-ray analyser. The average grain size of the consolidated samples was determined from the micrographs by applying the linear intercept method. The room temperature microhardness was measured on metallographically polished samples under a load of 1 N. Data points represent the average of 40 measurements each. Compression and tensile tests were carried out on spark-eroded and ground specimens. For compression testing, cylindrical specimens with a diameter of 4 mm and a length of 8 mm were used. The Zwick 1478 universal testing machine was equipped with a three-zone furnace to ensure a constant and uniform heating of both the specimen and the pistons. In order to minimize the influence of oxidation, the compression tests were run in an atmosphere of flowing Ar with a purity of 99.998%. The compressive strain was recorded via 3 alumina rods, directly pressed against the specimen or the upper die, respectively, and connected to a LVDT outside the heating zone of the furnace. The maximum
deformation detectable was limited to 1 mm, corresponding to a strain of 12.5%. Whereas compression tests were run with constant velocity, tensile tests on a Zwick 1484 testing machine were performed using constant strain rates. The corresponding specimens had a gauge length of 22 mm. Tensile tests could be carried out only in air.
3. Results 3.1. Microstructure The intermetallic/ceramic compounds, produced by high-energy-milling and HIP, are very homogeneous and consist of equi-axed g-TiAl-grains and x-Ti5(Si,Al)3particles. Depending on the adjusted volume fraction of the silicide phase, two types of microstructures may be distinguished. (a) Intermetallic/ceramic composites with a moderate silicide content of about 9.6 vol.% (as in Ti–45Al–2.4Si) or 13.1 vol.% (as in Ti–46Al–5Si) reveal a bimodal grain size distribution, i.e. submicron grains of the majority phase g-TiAl are surrounded by clearly smaller particles of the x-Ti5(Si,Al)3-phase (Fig. 1a). Different HIP-temperatures do not influence this overall appearance. Volume fractions and arrangement of the phases involved do not alter, only the phase specific mean grain sizes vary, but without changing their mutual ratio. An estimation of the finest grain size achievable under the applied HIP conditions is given in Ref. [18]. (b) Highly silicide containing compounds like Ti– 36Al–10Si, with about 30 vol.% of x-Ti5(Si,Al)3, do not show any difference between the mean grain sizes of the g-TiAl phase and the silicide phase (Fig. 1b). The silicide content is however still low enough to prevent the formation of a continuous ceramic network, i.e. the Ti5(Si,Al)3-grains may be regarded as isolated inclusions between the g-TiAl crystallites. Silicon-free binary TiAl-compounds reveal a globular microstructure of the g phase with small 2 grains precipitated at the triple points of g grain boundaries [19]. These alloys were included in the study as a submicrongrained, but silicon-free reference material. The microstructural data of all samples investigated were recently published in [4,16] along with their impurity contents. Considering the multistep powdermetallurgical processing of the samples, it should be noted that the amount of contamination by interstitial atoms like O (0.3 at.%) and N (0.1 at.%) is very low. Nevertheless, the material does not reach the purity of ingots. As the content of O and N probably exceeds the solubility limit of the g-TiAl phase, small amounts of oxides and nitrides should be present within the microstructure. A more detailed description of the microstructure is included in Refs. [16,17].
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Fig. 1. TEM micrographs showing the microstructure of TiAlSi compounds in the as-HIP’ed condition; (a) Ti–45Al–2.4Si, consolidated at 875 C; (b) Ti–36Al–10Si, compacted at 1000 C. Further details are provided in Ref. [16].
3.2. Temperature dependence of yield stress At low temperatures, submicron-grained alloys based on g-TiAl are characterized by high values of hardness and compressive yield strength [20]. As previously reported [4], this behavior can be explained by the clasp sical Hall–Petch relationship, resulting in a 1/ ddependence of grain size d on yield strength y , i.e. dislocation glide and mechanical twinning as the prevailing deformation mechanisms are hampered by the finemeshed net of grain-boundaries present in submicrongrained alloys. Consequently, further grain refinement leads to a continued increase of hardness and strength. In the high temperature range above 500 C, these relations change completely, i.e. ultrafine-grained materials become very soft. As demonstrated in Fig. 2, there is a
Fig. 2. Flow stress after 1.25% plastic compression strain, 1.25, in dependence of the deformation temperature for silicon-free alloys; ML, AM and CF refer to Mechanically Alloyed [19], high energy Attrition Milled (this work) and Cast/Forged material [21].
grain size-dependent inversion of yield strength, which means that the smaller the grain size of the specimen, the stronger and sharper the drop in yield strength upon raising the temperature. Silicon-doped g-TiAl-based compounds with a fraction of up to 32 vol.% Ti5(Si,Al)3 show a similar behavior as the binary TiAl alloys (Fig. 3). Grain refinement causes a reduction of flow stress, suggesting favorable conditions for hot-working of ultrafine-grained TiAlSi compounds. 3.3. Strain hardening Increasing the test temperature leads to a reduction of strain hardening. As demonstrated in Fig. 4, a binary TiAl sample with a mean grain size of 670 nm deforms already at 700 C under constant stress, i.e. the dynamic sample loading with a given velocity of the crosshead is comparable to a static loading under a corresponding constant stress. Thus, the dynamic compression tests are
Fig. 3. Flow stress 1.25 vs. temperature for Ti–46Al–5Si and Ti– 36Al–10Si with a silicide fraction of 13 or 32 vol.%, respectively.
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Fig. 4. Compression curves obtained on specimens of composition Ti– 48.9Al with a mean grain size of 670 nm.
formally quite similar to the steady state regime of creep tests.1 Specimens with smaller grain sizes show a lack of strain hardening at still lower temperatures. 3.4. Strain rate dependence of flow stress Submicron-grained TiAlSi compounds are characterized by a marked strain rate sensitivity at temperatures 5500 C. In Fig. 5, the flow stress of differently finegrained specimens of composition Ti–45Al–2.4Si is depicted as a function of the applied compression rate at a temperature of 800 C. The coefficient of strain rate sensitivity, m, that describes the slope of the respective curves, varies between 0.3 and 0.55. Values of that magnitude are indicative of superplastic deformation behavior. However, in order to actually demonstrate the potential of submicron-grained g-TiAl-based alloys for superplastic deformation at comparatively low temperatures, tensile tests are required. The corresponding results are presented in Section 3.6. The strong strain rate dependence of submicrongrained TiAlSi compounds has to be kept in mind if the temperature dependence of the flow stress (Figs. 2 and 3) is interpreted. Especially, from a high flow stress (under the strain rate conditions applied), it does not automatically follow that a good creep resistance will be observed. At constant stress, decreasing the grain size allows for higher strain rates to be adjusted in order to reach the maximum strain rate sensitivity. In the present case, the bisection of the mean grain size admits about a fivefold increase of the strain rate.
1 The expression ‘‘creep test’’ may be somewhat irritating, particularly since deformation rates in the range of 104–103 s1 are considered. However, it should be pointed out that similar deformation curves as in Fig. 4 are expected by applying a constant stress instead of a constant crosshead velocity.
Fig. 5. Strain rate dependence of flow stress 3.0 (stress after 3.0% true plastic compression strain) at 800 C for compounds of Ti–45Al– 2.4Si with different mean grain sizes. Fig. 4 represents a completion of Fig. 5 in Ref. [13].
Similar diagrams as in Fig. 5 are attained if the grain size is kept constant, but the test temperature is changed. From Fig. 4 in Ref. [13], it is obvious that grain refinement is comparable to a temperature increase with respect to the deformation behavior. This fact is of great technological interest, as it allows hot-working of submicron-grained TiAl alloys at reduced temperatures. 3.5. Microstructural stability Nano- and submicron-grained microstructures are prone to grain growth due to the large amount of stored grain boundary enthalpy. Therefore, it has to be investigated to what extent the mechanical data presented in the previous sections still refer to the as-HIP’ed microstructures shown in Fig. 1. The grain size of the test specimens was controlled twice. Firstly, some samples were annealed at the deformation temperatures of interest and subsequently cooled down to room temperature. The dwell time chosen for these heat treatments was 5 h, which is about 2 h more than the preheating time necessary to adjust a steady temperature profile within the ceramic pistons of the testing machine. This gives a reference for the microstructure at the start of the deformation procedure. Secondly, the deformed specimens were investigated in order to get informations about microstructural changes occurring upon deformation. As mentioned before, room temperature hardness and yield strength of the material are closely related to the grain size via the Hall-Petch correlation. Therefore, instead of the time-consuming preparation of TEM foils, it was sufficient to perform hardness measurements in order to get a first impression of microstructural changes. Occasional spot checks carried out in the TEM revealed that these microstructural changes were in fact confined to grain growth processes. The
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Fig. 7. Room temperature hardness, obtained on specimens of composition Ti–45Al–2.4Si after 12.5% compressive deformation at 800 C, in dependence of the applied strain rate. Hardness data referring to tensile strained specimens ‘‘B’’ and ‘‘F’’ is also included (see Section 3.6 for more details).
Fig. 6. (a) Room temperature hardness as a function of annealing temperature for three g-TiAl based compounds with different fractions of the silicide phase Ti5(Si,Al)3, according to 0!10!13 vol.%. The resulting grain growth at temperatures corresponding to the applied HIP-temperatures is calculated on basis of the Hall–Petch parameters given in Ref. [4]. (b) Temperature-dependent flow stress 1.25 of the same compounds.
results of investigations on the microstructural stability are shown in Figs. 6 and 7. In Fig. 6a, the hardness at room temperature is plotted against the chosen annealing temperature. As long as a high hardness level is retained, grain growth does not occur. The applied HIPtemperatures are also indicated. As expected, the stability of the microstructure at the respective HIP-temperature is enhanced by increasing the amount of the silicide phase (0!10!13 vol.%). In Fig. 6b, the yield stresses of the different compounds are shown as a function of the deformation temperature. The three curves widely overlap. In combination, Fig. 6a and b show that the marked drop of yield strength occurs already before the onset of considerable grain growth, at least in the case of the silicon-doped material. Accordingly, the decrease of flow stress is rather caused by a change of the deformation mode than by microstructural instabilities of the test specimens. This is an important result with respect to technological applications: The low flow stress due to the fine grain size and the silicide-dispersion-induced stability of the microstructure admit favorable hot-working conditions for the shaping of components.
After compression straining at 800 C, specimens of alloy type Ti–45Al–2.4Si do not show any difference in hardness between the unstrained or the deformed state, independent of the strain rate applied (Fig. 7). This allows the conclusion, that the initial microstructure and the respective mechanical properties did not significantly change. However, it also has to be kept in mind that the total compressive strain amounts to only 12.5%, which might be not enough for the start of strain-induced reactions like recrystallization. Regarding tensile tests, the preconditions are somewhat different, as discussed in Section 4. 3.6. Tensile straining Tensile tests at 800 C were performed on the basis of the mechanical data derived from compression straining of Ti–45Al–2.4Si. The experiments were launched to validate the supposed potential for superplastic deformability and to verify the results obtained by compressive loading. The results are summarized in Table 1, and may be described as follows: : 1. At strain rates "43:2103 s1 elongations "5175% were reached (Fig. 8). Larger deformations could not be managed due to the limited length of the : furnace. By increasing the strain rate to "46:4 103 s1 the specimens fail after elongations of 103 or 113%, respectively. 2. The coefficients of the strain rate sensitivity are comparable to those obtained by compression testing. Increasing the strain rate leads to higher strain rate sensitivities. : 3. Simultaneously to an increase of " and m, the fracture strain drops. However, necking of the samples is not observed.
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Table 1 : Coefficient of strain rate sensitivity, fracture strain, porosity and resulting grain size of Ti–45Al–2.4Si specimens deformed at different strain rates. "1 : and "2 are the strain rates applied during the strain rate cycle tests
Fig. 9. TEM micrograph, taken from the neck region of specimen ‘‘B’’ after a deformation of 175%. While the grain shape did not change, slight grain coarsening from 170!248 nm can be observed.
Fig. 8. Tensile samples of the compound Ti–45Al–2.4 Si after deformation at 800 C in air. An undeformed specimen is shown for comparison.
4. The microstructure has conserved its equi-axed grain character after deformation (Fig. 9). Contrary to compression loaded specimens, limited grain growth occurs. However, within the strain rate range investigated, increasing the deformation rate leads to a reduction of grain growth. 5. Measurements of the room temperature hardness were carried out in both the neck and the head region of deformed specimens. The results are presented in Table 1 and Fig. 7. Whereas the hardness of the specimen heads did not show any significant changes in comparison to
the as-HIP’ed state, the neck regions have become considerably softer, irrespective of the strain rate applied. As mentioned above, grain growth proved to depend on the strain rate and thus can not explain the loss of hardness. Instead, the softening is attributed to the formation of voids, especially in fast strained specimens (Fig. 10). The accumulation of pores near surface notches indicates that local stress concentrations largely contribute to cavitation.
4. Discussion The high temperature mechanical behavior of submicron-grained g-TiAl-based compounds is essentially characterized by the strongly reduced flow stress. This property already allowed for rather low consolidation temperatures during powder processing and implicitly
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napp
Fig. 10. Evolution of voids in specimen ‘‘F’’ upon tensile straining at 800 C. The SEM micrograph is taken at the edge of the central neck region of tensile specimen ‘‘F’’. The surface roughness leads to a nonuniform distribution of voids. A lot of pores have accumulated at the tip of a notch (arrowed), whereas other regions are essentially free of voids (encircled).
contributed to the fine-grained microstructure. In order to elucidate the mechanisms of high temperature deformation, the stress-strain curves will be further analyzed. As indicated above, the compression tests can be treated like static deformation tests, i.e. the compressive : strain rate " — which is actually a given parameter — was analysed in dependence of the plateau stress 3:0 arising after an initial plastic deformation of three per: cent. Generally, " is a function of temperature T, stress and quite a lot of microstructural parameters Si , like grain size, grain shape, volume fractions and distribution of phases involved etc.: : " ¼ fðT; ; Si Þ
ð1Þ
In the simplest case, these parameters do not interact. Assuming that essentially one process is dominating, the analysis may be based on the equation n Q : " ¼ A
exp E RT
ð2Þ
E denotes the Young modulus, n is the stress exponent and Q refers to the activation energy of the deformation step that is rate-controlling. The influence of the microstructural parameters is represented by the prefactor A. The slope of the curves shown in Fig. 5 leads to the respective coefficient of the strain rate sensitivity, m, which is reciprocal to the apparent stress exponent
: dln" 1 ¼ ¼ dln T¼const: m
565
ð3Þ
For small stresses, napp is about equal to n as long as the prefactor A does not depend on stress. At larger stresses, napp exceeds n. In advance to the analysis, several aspects concerning the microstructural stability upon high temperature deformation have to be assessed. 1. The mean grain size before and after deformation was evaluated by hardness measurements. However, this procedure only offers a selective shot of the truth. For example, it is possible that a nonvariant grain size is the result of both grain growth and dynamic recrystallization, that balance each other. This assumption is supported by studies of Imayev et al. [22] and Koeppe et al. [23], suggesting that grain refinement or grain growth prevails depending on whether a certain strain rate or stress level is exceeded or not. The results of the tensile tests performed within this study tend to support this view. For the Ti–45Al–2.4Si compound with a mean grain size of 170 nm, this neutral strain rate is obviously larger than 3.2103 s1 at 800 C. Thus, it may be explained why the slowly strained specimen ‘‘B’’ shows more intense grain coarsening than specimen ‘‘F’’. In the case of the compression tests, it may be argued that the strain level attained upon deformation should be too low to observe similar effects. 2. In strain rate cycle tests performed both in tension and under compression, the doubling of the strain rate does not lead to a spontaneous change of flow stress. Instead, the corresponding stress level is reached rather gradually, which is indicative of microstructural changes occurring inside the specimen. Apart from variations of the grain size, the number and arrangement of dislocations, twins etc. could account for this behavior. Consequently, the structure-sensitive prefactor A in Eq. (2) should not virtually be treated as a constant, and napp should deviate from n. However, the actual differ: ence must be very small, as strain rate changes from " i : to " iþ1 (and back) lead to the same plateau stresses that are reached if the deformation tests are performed with : a strain rate "iþ1 right from the beginning. Summarizing, the question of microstructural stability can be regarded as rather uncritical with respect to the performed analysis. In this study, napp amounts to values between 1.8 and 3.3. Superplasticity based on grain boundary sliding requires n=2 [24]. The unchanged grain shape after a tensile strain of 175% provides strong evidence for grain boundary sliding to play an important role. In contrast to this, deformation modes based on diffusional flow like Coble [11] or Nabarro– Herring creep [25] are characterized by n ¼ 1 and would lead to an elongation of the crystallites. As this is not observed, it is concluded that napp does not significantly overestimate n.
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Grain boundary sliding as an essential part of the superplastic deformation must always be supported by accommodation processes that prevent the formation of voids at grain triple junctions. The temperature dependence of deformation can provide further information about the nature of the assisting mechanism. By plotting : " versus 1/T, the apparent activation energy may be derived according to 0 1 : B dln" C Qapp ¼ R @ ð4Þ A 1 dln T ¼const:
small particles require higher stresses for dislocation generation and multiplication. More information about the deformation mode can be drawn from the grain size dependence of the strain rate. The microstructure-sensitive parameter A may be expressed in the form A / dp
ð5Þ
In Fig. 11, the modulus-compensated strain rate under compression is plotted against 1/T. The slope of the curves leads to a value of Qapp ¼ 351 kJ/mol, irrespective of whether the compounds Ti–45Al–2.4Si or Ti– 36Al–10Si are regarded. Obviously, the same mechanisms are operative in both materials. This is not surprising, assuming that grain boundary sliding is assisted by processes controlled by only the g-TiAl phase. Though Frommeyer et al. specify an apparent activation energy of 350 kJ/mol also for the deformation of monolythic Ti5Si3 at temperatures between 1000 and 1300 C [27], it seems unlikely that the hard x-Ti5(Si,Al)3 phase is codeformed within a composite containing the softer gTiAl phase. This assumption is substantiated by studies of Wang et al. on powder-metallurgically processed specimens of composition Ti–45Al–2.7Si, showing that the silicide particles are free of dislocations after deformation at 700 C [28]. The same should be valid with respect to the still finer silicide grains of this study, especially as
The grain size exponent p results from the slope of a : ln" vs. lnd plot (Fig. 12). At T ¼ 800 C and ¼ 150 MPa, a grain size exponent of p ¼ 2:2 is obtained. Classifying the experimentally derived parameters, Qapp , napp and p the deformation mechanism may be characterized as follows: The apparent activation energy of 351 kJ/mol is somewhat higher than the reported activation energy of 250–290 kJ/mol for the titanium self diffusion in g-TiAl [29,30]. Nevertheless, it compares well to the published activation energies of comparatively coarse-grained g-TiAl alloys (Table 2), that were deformed at similar or even higher (>1000 C) temperatures. In connection with a grain size parameter closer to 2, it is concluded that processes based rather on volume diffusion than on grain boundary diffusion support the grain boundary sliding of submicrocrystalline TiAlSi compounds. A corresponding model is sketched in Fig. 13. The diffusional flow inside the g-TiAl grains may be provided by dislocation climb, for example. The stress and grain size exponents are very similar to the values published by Ameyama et al. [14] and Mishra et al. [15] for the superplastic deformation of gTiAl-based alloys with grain sizes of about 1 mm [14] or 20 mm [15], respectively. Thus, the reduction of the
Fig. 11. True compressive strain rate, compensated for the temperature dependence [26] of the Young’s modulus, as a function of the reciprocal temperature for compounds of a type Ti–45Al–2.4Si and Ti–36Al–10Si. The analysis was performed at a stress level of 150 MPa.
Fig. 12. True compressive strain rate in dependence of the mean grain size, demonstrated for samples of composition Ti–45Al–2.4Si. The grain size exponent of the deformation was derived for a temperature of 800 C and a stress level of 150 MPa.
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Table 2 Stress exponents and apparent activation energies obtained from high temperature deformation experiments performed on g-TiAl based alloys [14–15], [19], [31–37]. a and b refer to tensile or compression tests, respectively. The alloys marked by * could be superplastically deformed under the conditions specified
mean grain size down to the low submicron range does not cause a change in the prevailing deformation mechanism, it rather decreases the lower boundary of the temperature regime where these mechanisms work. The results of this study are in some contrast to recently published work of Imayev et al. [36]. According to a specified activation energy of 180–200 kJ/mol for the superplastic deformation of submicron-grained TiAl-based alloys at moderate temperatures of about 750–850 C, the authors claim deformation to be accomplished by grain boundary sliding assisted by dif-
fusion along the grain boundaries. The reason for this discrepancy cannot be ascertained. However, it should be noted, that in a corresponding temperature range, similar specimens as those investigated in [36] require an activation energy of 335 kJ/mol for grain growth [39], which rather favours volume instead of grain-boundary diffusion to play a dominant role. The formation of voids upon superplastic deformation of g-TiAl alloys is a frequently observed phenomenon [40–42]. Reducing the grain size was shown to be an appropriate countermeasure [41]. In the present case, a
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Fig. 13. Model for the superplastic deformation of submicron-grained Ti–45Al–2.4Si: grain-boundary sliding, accommodated by dislocation creep inside the g-TiAl crystals (picture taken from Refs. [38]).
lot of pores are visible in the vicinity of surface notches. Thus it is concluded that stress-induced cavitation plays an important role. On a submicron scale, hard silicide particles are supposed to be responsible for the occurrence of local stress concentrations leading to the formation of voids all over the neck of the sample. This may be illustrated by replacing grain 3 of Fig. 13 by a silicide particle impermeable for dislocations. In general, the grain triple points can be identified as preferential nucleation sites for the formation of cracks [43]. According to Stroh [44], cracks may be initiated if the shear stress exceeds a critical value of c ¼
12G L
ð6Þ
with denoting the surface energy and L representing the length of the glide area. From Eq. (6) it is obvious that grain refinement allows for higher critical stresses. On the other hand, a lot of triple points are occupied by silicide particles that could reduce the surface energy and thus balance out the positive influence of a small grain size. In this respect, an even finer distribution of smaller silicides might be useful.
5. Summary and conclusions 1. Multiphase g-TiAl-based compounds with a very fine-grained microstructure were produced by high energy milling and hot isostatic pressing. The grain size of the g-TiAl matrix amounts to about 160–480 nm,
whereas the size of precipitated x-Ti5(Si,Al)3 particles varies between 80 and 190, depending on the HIP conditions. 2. At temperatures above 500 C, these alloys become very soft. The smaller the grain size, the stronger the drop of flow stress. Deformation occurs without any strain hardening. 3. During deformation at high temperatures, the finely dispersed silicide particles prevent grain growth as long as the HIP temperature is not exceeded. 4. The flow stress is rather sensitive to the applied strain rate. Superplasticity becomes feasible at 800 C, allowing elongations 5175%. The apparent activation energy of deformation was assessed to be Qapp ¼ 351 kJ/ mol. In connection with a stress exponent around n ¼ 2 and a grain size exponent of p ¼ 2:2, it is concluded that the superplastic deformation is accomplished by grain boundary sliding accommodated by diffusional processes inside the g-TiAl grains. Such a mechanism does not differ from the deformation modes described for conventionally grained materials deformed at temperatures 51000 C. 5. The silicide phase is not directly involved in the process of deformation. In particular, Ti–36Al–10Si compounds with about 32 vol.% of the silicide phase reveal the same activation energy as alloys of composition Ti–45Al–2.4 Si with only 9.6 vol.% Ti5(Si,Al)3. In both cases only the g-TiAl phase seems to contribute to the overall deformation. However, changes are expected if the silicide content is high enough to form a continuous network within the g-phase. 6. Increasing the tensile strain rate supports the formation of stress induced voids that reduce the overall deformability. The silicide particles are supposed to promote cavitation.
Acknowledgements This work has been supported by the Deutsche Forschungsgemeinschaft (German Science Foundation) within the scope of Sonderforschungsbereich 371. Thanks are due to Dr. R. Gerling for providing the powders for high energy milling. The experimental help of U. Lorenz and R. Behn is gratefully acknowledged.
References [1] Gertsman VY, Hoffmann M, Gleiter H, Birringer R. Acta Metall Mater 1994;42:3539. [2] Sanders PG, Eastman JA, Weertman JR. Acta Mater 1997; 45:4019. [3] Koch CC, Morris DG, Lu K, Inoue A. MRS Bulletin 1999;24:54. [4] Bohn R, Klassen T, Bormann R. Acta Mater 2001;49:299. [5] Wang DL, Kong QP, Shui JP. Scripta Metall Mater 1994;3:47. [6] Deng J, Wang DL, Kong QP, Shui JP. Scripta Metall Mater 1995;32:349.
R. Bohn et al. / Intermetallics 9 (2001) 559–569 [7] Xiao M, Kong QP. Scripta Metall Mater 1997;36:299. [8] Sanders PG, Rittner M, Kiedaisch E, Weertman JR, Kung H, Lu YC. Nanostr Mater 1997;9:433. [9] Hahn H, Averback RS. Nanostr Mater 1992;1:95. [10] McMeeking RM, Kuhn LT. Acta Mater 1992;40:961. [11] Coble RL. J Appl Phys 1963;34:1679. [12] Bohn R, Fanta G, Bormann R. In: Morris DG et al., editors. Intermetallics and superalloys. Weinheim, Germany: WileyVCH, 2000. p. 271. [13] Fanta G, Bohn R, Dahms M, Klassen T, Bormann R. Intermetallics 2001;9:45. [14] Ameyama K, Uno H, Tokizane M. Intermetallics 1994;2:315. [15] Mishra RS, Lee WB, Mukherjee AK, Kim Y-W. In: Kim Y-W et al., editors. Gamma titanium aluminides. Warrendale, USA: TMS, 1995. p. 571. [16] Bohn R, Fanta G, Klassen T, Bormann R. J Mater Res, in press. [17] Bohn R, PhD thesis, Technische Universita¨t Hamburg-Harburg, 1999 (in German). [18] Bohn R, Klassen T, Bormann R. Adv Eng Mater 2001;3:238. [19] Oehring M, Appel F, Pfullmann Th, Bormann R. Appl Phys Lett 1995;66:941. [20] Bohn R, Oehring M, Pfullmann T, Appel F, Bormann R. In: Suryanarayana C et al., editors. Processing and Properties of Nanocrystalline Materials. Warrendale, USA: TMS, 1996. p. 355. [21] Appel F, Lorenz U, Sparka U, Wagner R. In: Oikawa et al., editors. Strength of metals — Proc. of the 10th Int. Conference ICSMA 10. Sendai, Japan: JIM, 1994. p. 341. [22] Imayev RM, Imayev VM, Salishchev GA. J Mater Sci 1992; 27:4465. [23] Koeppe C, Bartels A, Clemens H, Schretter P, Glatz W. Mater Sci Eng 1995;A201:182. [24] Lu¨thy H, White RA, Sherby OD. Mater Sci Eng 1979;39:211. [25] Herring C. J Appl Phys 1950;21:437. [26] Schafrik RE. Metall Trans 1977;8A:1003.
569
[27] Frommeyer G, Rosenkranz R, Lu¨decke CZ. Metallkd 1990; 8:307. [28] Wang G-X, Dogan B, Hsu F-Y, Klaar H-J, Dahms M. Metall Trans 1995;26A:691. [29] Mishin Y, Herzig C. Acta Mater 2000;48:589. [30] Kroll S, Mehrer H, Stolwijk N, Herzig C, Rosenkranz R, Frommeyer GZ. Metallkd 1992;83:591. [31] Seetharaman V, Lombard CM. In: Kim Y-W et al., editors. Microstructure/property relationships in titanium aluminides and alloys. Warrendale, USA: TMS, 1991. p. 237. [32] Martin PL, Mendiratta MG, Lipsitt HA. Metall Trans A 1983; 14A:2170. [33] Kampe SL, Bryant JD, Christodoulou L. Metall Trans A 1991; 22A:447. [34] Cheng SC, Wolfenstine J, Sherby OD. Metall Trans A 1992; 23A:1509. [35] Wang JG, Nieh TG. Intermetallics 2000;8:737. [36] Imayev RM, Salishchev GA, Senkov ON, Imayev VM, Shagiev MR, Gabdullin NK, Kuznetsov AV, Froes FH. Mater Sci Eng A 2001;A300:263. [37] Rommerskirchen M. PhD thesis, RWTH Aachen, 1997 (in German). [38] Sherby OD, Wadsworth J. Progr Mater Sci 1989;33:169. [39] Senkov ON, O¨vec¸oglu ML, Srisukhumbowornchai N, Froes FH. Nanostr Mater 1998;10:935. [40] Lee WB, Yang HS, Kim Y-W, Mukherjee AK. Scripta Metall Mater 1993;29:1403. [41] Imayev VM, Imayev RM, Salishchev GA. Mater Sci Eng 1996; A208:226. [42] Appel F, Clemens H, Glatz W, Wagner R. MRS Symp Proc 1997;460:195. [43] Edington JW, Melton KN, Cutler CP. Prog Mater Sci 1976; 21(2):61. [44] Stroh AN. Proc Roy Soc 1934;A223:404.