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Microstructure and plasticity of two molybdenum-base alloys (TZM)
Materials Science and Engineering, A 160 (1993) 189-199
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Microstructure and plasticity of two molybdenum-base alloys (TZM) H. A. Calderon and G. Kostorz Institut fiir Angewandte Physik, E TH Ziirich, CH-8093 Ziirich (Switzerland)
G. Ullrich Paul-Scherrer-Institut, CH-5232 Villigen PSI (Switzerland) (ReceivedMay 10, 1991; in revisedform September22, 1992)
Abstract In two different commercial Mo-base alloys (TZM), produced by vacuum melting and by powder metallurgy respectively, microstructural differences (particle size and chemistry, grain and subgrain sizes, dislocation density) have been found which affect the observed mechanical properties of the material. The compression-creep properties at 1423 K show a negligibly small creep rate at a stress of approximately 200 MPa. Trapping of dislocations by particles is proposed to be the controlling deformation mechanism during creep. The microstructure and fatigue properties of TZM welds were also investigated. Friction welds showed the best mechanical properties. Fatigue measurements in load control at room temperature and 1123 K show that the endurance limit of the vacuum-melted alloy is higher than that of the powdermetallurgically processed alloy.
1. Introduction Mo-base alloys are commonly used for applications at high temperatures where a high mechanical strength is required. The creep strength of Mo-base alloys at temperatures higher than l l 0 0 K is considerably better than that of Ni alloys [1]. Nevertheless, the poor oxidation resistance at temperatures higher than 750 K limits the use of these alloys to well controlled environments. Among the Mo alloys of technological interest, the T Z M alloy is the most important. Its applications range from turbine blades to high-temperature forging dies. The known good mechanical properties of TZM also made it attractive for application in a project involving high temperature helium turbines [2, 3]. The name of this alloy is related to its composition; Mo-0.5Ti-0.08Zr-(0.1-0.4)C (in atomic per cent). The presence of C leads to the formation of a hardening dispersion of carbides which strengthens the alloy beyond the limits shown by single-phase Mo alloys, i.e. solid-solution hardened Mo alloys containing W or Re [4]. T Z M is commercially produced via two industrial processes, vacuum melting (VM-TZM) and powder metallurgy (PM-TZM). The latter fabrication method is considerably less expensive but some reports [5, 6] indicate that the creep-rupture properties of PM-TZM are inferior to those of VM-TZM. In extensive investigations of the creep properties of PM-TZM conducted by Rosen and coworkers [7], however, the effects of 0921-5093/93/$6.00
particle dispersion and thermomechanical treatment on the creep properties were studied, and an excellent creep strength was found. It was concluded that the creep strength stems from a preferential nucleation of precipitates at subgrain boundaries. However, most of the details of the relevant hardening mechanisms and of the microstructure of this alloy are still unknown. In the present work, some creep properties of the two different T Z M alloys were studied. Furthermore, the relevant microstructures of both alloys were compared in the as-received state and after creep. From these observations the mechanism likely to control the creep rate may be identified. Several investigations were also conducted in order to study the conditions under which T Z M can be welded, with suitable mechanical properties of the welds (see for example refs. 8 and 9). In general, most industrial processes produce welds with inferior mechanical properties. This is due to recrystallization and pore formation induced by the welding process [8, 9]. Some fatigue properties of welded and base materials are reported. Finally, the microhardness and the microstructures produced by two welding procedures are compared.
2. Experimental procedures The specimens for the study of the creep properties in uniaxial compression were obtained from bars © 1993 - Elsevier Sequoia. All fights reserved
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50 mm in diameter by cutting pieces of a square crosssection with an aspect ratio lid of approximately 2 (here I is the length of the specimen and d the equivalent diameter), in order to minimize buckling and barreling. A thin layer of BN was used as a lubricant. All tests were performed in a screw-driven machine programmed to keep the applied stress constant. The testing temperature was 1423 K, and a vacuum better than 5 x 10-3 Pa was maintained during the tests. Transmission electron microscopy (TEM) was performed in a JEOL 100C microscope. Thin foils were produced by electropolishing in a twin-jet apparatus with a solution of 12.5 vol.% H2SO 4 in ethanol at 258 K. A voltage of 10 V produced a current density of approximately 3 mA mm- 2. Microstructures after welding were observed after either electron beam welding two plates of a thickness of 8 mm or friction welding bars of a diameter of 25 mm. The specimens were stress relieved at 973 K f o r l h. The fatigue tests were performed in a screw-driven machine (Schenk, Type PH 1G 0016) in load control at a frequency of 25 Hz (sinusoidal waveform), with a ratio O'max/Ormi n = - 1. The fatigue specimens had a cylindrical gage section with a diameter of 1.7 mm or 2 mm. The gage length was five times the diameter. The high-temperature (T = 1123 K) tests took place in a helium atmosphere that prevented the production of volatile Mo oxides. This atmosphere simulated the environment of a high-temperature gas-cooled reactor. It was also possible to add controlled amounts of impurities to the helium atmosphere.
Two Mo-base alloys TABLE 1. Grain and subgrain sizes in TZM alloys Location
PM-TZM l (/~m) (/~m)
Grain size Edge of bar 38 Center of bar 33 Subgrain size Edge of bar Center of bar
1.2 1
114 128
VM-TZM R=lgl¢v ¢v l (#m) (/~m)
R=~I~
3 3.9
14.9 15.6
0.7 1.7 0.74 1.4
19 26 3.6 5.1
284 406 0.9 0.8
4 6.4
3. Experimental results 3.1. Microstructural observations 3.1.1. As-received specimens The two TZM alloys have different grain structures, as summarized in Table 1. The grains of PM-TZM are elongated parallel to the extrusion direction. The aspect ratio ]g/~ (/g and ~ represent the mean length and thickness respectively) is relatively low. The aspect ratio for VM-TZM is considerably higher than for PMTZM. In PM-TZM, coarse particles can be seen as shown in Fig. l(a), their sizes ranging from 1 to 15/am. Some of these particles are cracked and exhibit poor bonding to the matrix. Energy-dispersive X-ray fluorescence analysis shows that these large particles contain mostly Zr with minor amounts of Ti and Ca. Electron diffraction reveals a cubic structure of the particles, with a lattice parameter of around 5.17 A. Therefore, it is very likely that the particles are ZrO2 which could be stabilized at room temperature by Ca. The grain struc-
Fig. 1. Grain microstructuresof the as-received samples, parallel to the extrusion direction: (a) PM-TZM (the dark, coarse particles are ZrO2); (b) VM-TZM.
ture of VM-TZM is typical of cold-worked materials (Fig. l(b)). No coarse particles are found in this case. TEM also reveals considerable differences in the microstructure of the two as-received TZM alloys. The subgrain size and aspect ratio, the dislocation arrangement and the carbide distribution and size are all characteristic of each alloy. Figure 2 shows typical particle populations in the two alloys. The particles in PM-TZM (Fig. 2(a)) are in general homogeneously distributed in subgrain boundaries and in the subgrain
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interiors. The particle contrast suggests that precipitates are coherent or semicoherent and spherical. Considering the alloy composition, and the fact that most of the Zr forms ZrO 2 (Fig. l(a)), the small particles in Fig. 2(a) should be TiC. This is supported by TEM observations of the strain contrast expected for such particles. However, the average particle sizes are not independent of the specimen location in the 50 mm bar. A summary of the measurements is given in Table 2. In VM-TZM, precipitate size and morphology are considerably different from those of PM-TZM (Fig. 2(b)). The particles are not perfectly spherical and are apparently incoherent or only partially coherent with the matrix. Strain contrast is observed only with highorder reflections. The dislocation structure in this alloy shows a higher degree of recovery. Some dislocation networks can be observed in an advanced degree of development. Furthermore, many particles are connected by dislocations which probably promotes accelerated coarsening of particles. Table 2 summarizes the metallographic measurements. It is important to note that particle sizes and dislocation densities
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are not constant across the diameter of the original bar, which can be linked to the observed variation in mechanical behavior (see below). In VM-TZM, no large oxide particles are found and therefore the particle chemistry should be ZrxTi,_xC. The lattice mismatch complete coherency, thus reducing between Mo and ZrxTi,_~C is large enough to prevent strain contrast in TEM for low-order reflections. The exact amount of Ti in these carbides was not determined experimentally. However, measurements of lattice parameters in TEM and comparison with published data [10] indicate that x is approximately 0.5-0.6. The volume fraction fv of precipitates can be calculated from the nominal composition to range from 0.24% to 0.36%. Experimentally, fv = 1%-1.5% was found. 3.1.2. Welds Microhardness measurements. Figures 3 and 4 show
the microhardness profiles through the welded areas in the two alloys after electron-beam and friction welding. All measurements were taken in a direction perpendicular to the weld. In both alloys the microhardness values of the welded regions and of the heat-affected zones remain constant or show little change after friction welding. In contrast, electron-beam welding causes
TABLE 2. Dislocation density and average particle sizes in TZM alloys Location
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VM-TZM
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46.3+ 1 33.1 + 1
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Fig. 3. Microhardness (HV) profiles taken across the welds of PM-TZM: x electron-beam weld; • friction weld.
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a decrease in microhardness by about 30% in both TZM alloys. The quasi-homogeneous microhardness observed for the friction-welded alloys indicates the relatively low damage induced by this welding technique, whereas a considerable reduction in strength is brought about by electron-beam welding. The absolute
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Two Mo-base alloys
values of microhardness depend on the thermomechanical history of the specimen and will not be discussed here.
Microstructure after electron-beam welding. The grain structure of the weld region in both TZM alloys shows large differences in grain size as shown in Fig. 5(a)-(c) for VM-TZM. In this figure, the weld (Fig. 5(a)), the heat-affected zone (HAZ) (Fig. 5(b)-(c)) and the base material (Fig. 5(c)) can be observed. Recrystallization in the weld and the H A Z is evident. Similar features can be seen in the case of PM-TZM but, in addition, there are large pores in the weld, with sizes ranging from 20/~m to 200/~m (Fig. 5(d)). The size of the H A Z is approximately 12 mm in both cases. Furthermore, transgranular cracks were found in the VM-TZM weld. This is probably due to an inappropriate welding atmosphere [8]. Table 3 shows the grain sizes measured in the weld and the HAZ. Figure 6 shows the grain microstructure as observed in the scanning electron microscope. The weld of PMTZM has large particles and pores inside the grains. The bright areas that decorate the grain and subgrain boundaries are most probably not an artefact. Subse-
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Fig. 5. Grain structure after electron-beam welding:(a) VM-TZM weld; (b) VM-TZMheat affected zone; (c) VM-TZMbase material; (d) PM-TZMweld.
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TABLE 3. Grain sizes in welds and heat affected zones (HAZ) Alloy
Processa
Region
Grain size (mm)
PM-TZM
EBW
Weld
PM-TZM VM-TZM
EBW EBW
HAZ Weld
VM-TZM PM-TZM PM-TZM VM-TZM VM-TZM
EBW FW FW FW FW
HAZ Weld HAZ Weld HAZ
0.25 width 0.80 length 0.70 0.25 width 0.75 length 0.060 0.008 0.025 0.007 0.030
aEBW electron beam welding, FW friction welding.
Fig. 6. Grain structure of PM-TZM after an electron-beam welding as observed by SEM. The particles at the grain and subgrain boundaries are presumably (Ti,Mo)2C. The arrowed particles are ZrO 2. quent TEM studies revealed a high particle density at grain boundaries. These particles are probably Mo carbides (see below). The larger particles marked by arrows are Zr20 as described previously. TEM shows that the dislocation density in the weld and the H A Z of PM-TZM is very low, typical of a recrystallized structure. The dislocation density is relatively high only around large particles. Very few semicoherent TiC particles are observed in the weld. Instead, large carbide particles rich in Ti and Mo are visible along the grain and subgrain boundaries. These particles have sizes ranging from 250 nm to 2/tm. Wadsworth et al. [9] have identified similar particles as (Ti,Mo)2C. The weld in VM-TZM also shows a low dislocation density (less than 107 cm-2). Only large, elongated (Ti,Mo)2C particles are observed, mostly along grain boundaries. The finer ZrxTil_xC particles found in the as-received alloy are not present. Figure 7(a) illustrates the grain boundary carbides, and Fig.
Fig. 7. (a) Grain boundary carbides in the weld region of VMTZM after friction welding; (b) corresponding diffraction pattern. 7(b) shows the corresponding diffraction pattern. The H A Z of the VM-TZM weld shows elongated subgrains, thus recrystallization did not occur in this case. Areas with large ZrxTil_xC particles are found. As in the as-received alloy, strong dislocation-particle inte"actions are observed. The large (Ti, Mo)2C particles are also observed along the grain boundaries, but less frequently than in the PM-TZM alloy. Microstructure after friction welding. The weld has a fine-grained structure in both alloys as shown in Fig. 8. The extension of the welds and H A Z s is smaller than 1.5 mm. This is considerably less than for alloys which are electron beam welded. In the PM-TZM weld, a
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relatively high number of particles (size 10-15/am) can be seen (Fig. 8(a)). These particles are rich in Ti and/or Mo, and most of them are carbides. Table 3 lists the measurements of grain size. TEM shows that the H A Z of PM-TZM has elongated subgrains (2-2.5 /am in length) with large Tiand/or Mo-rich carbides. Some signs of recrystallization can also be seen but this process is not complete, as indicated by frequently occurring areas of high dislocation density. Figure 9 shows representative views of the subgrain microstructure and the carbides in the weld. It can be observed (Fig. 9(b)) that hardening carbides (TIC) are still present within the welded region, increasing the mechanical strength of the weld in comparison with electron-beam welding. The subgrains in the VM-TZM weld are equiaxed with an average subgrain size of 2/am. In the weld, recrystallization is incipient, and areas with a high dislocation density can be observed. A TEM determination yields a dislocation density of 7 x 10 ~° cm -2 for these areas. This dislocation density is higher than that found in PM-TZM. Large carbides (1-2/am) are also, but infrequently, observed along subgrain boundaries. A few incoherent ZrxTil_xC particles are observed inter-
Fig. 8. Grain structure after friction welding: (a) PM-TZM (the large particles marked by arrows are M2C-typecarbides rich in Mo and Ti); (b) VM-TZM.
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Two Mo-base alloys
acting with and pinning dislocations. In the H A Z the subgrains vary more broadly in size, and some of them do not seem to be equiaxed. Very few large carbides are found. Instead, a fine homogeneous distribution of ZrxTi~_xC is observed. The mean size of these particles is 50 nm. These observations suggest that the grains and subgrains produced by the welding procedure are very fine and equiaxed and that no recrystallization takes place in these regions. Some deformation is also detectable. In both welds, large oxide particles are present. Analysis of diffraction patterns shows that they correspond to monoclinic MOO2. A dark-field view of these oxides in PM-TZM and the corresponding diffraction pattern with a zone axis close to [110] are shown in Fig. 10. The objective aperture used for the dark-field image (Fig. 10(b) was large enough to produce phase contrast with the transmitted beam and as MoO2 has large lattice constants, lattice fringes are visible for the (100) plane. In the HAZ, the incipient recrystallization produces a variable subgrain size and also precipitation of large carbide particles which are clearly deleterious for the
Fig. 9. Weld region of PM-TZM after friction welding: (a) subgrains and carbides;(b) TiC particles.
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mechanical properties of the weld. Hardening carbides (i.e. TiC for PM-TZM and TixZr ]_ xC in VM-TZM) are observed more frequently than in the weld.
3.2. Creep experiments 3.2.1. Creep rate Figures l l(a) and l l(b) show the stationary creep strain rate g at 1423 K as a function of the constant compressive stress used experimentally for the two alloys tested. In the figures, the positions of the specimens in the original bar are indicated by numbers, 5 representing the center and 1 the bar edge. Results for specimens 1 are not reported as the data scattered considerably. Figure 1 l(a) shows that there is a variation in the steady-state creep rate with specimen location for the VM-TZM bar. The peripheral region
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Fig. 11. Steady-state creep rate g vs. applied stress a in TZM alloys at 1423 K: (a) VM-TZM; (b) PM-TZM. The numbers indicate the specimen location in the original bar of a diameter of 50 mm; 1 indicates the bar edge and 5 the center.
shows a lower creep strength. In PM-TZM, the creep response is more uniform, and only small differences in g can be found. The data can be fitted to a power law oc a" with good data correlation. The stress exponent n takes the values 15 and 11 for the center and periphery of the PM-TZM bar and 14 for the center of the VM-TZM bar. At high stresses, the creep response of the central region of VM-TZM appears to be slightly better than that of the center of the PM-TZM bar. The steady-state creep rate induced by stresses lower than 180 MPa was below the resolution limit of the extensometer used.
3.2.2. Microstructure after creep deformation Representative views of the microstructure after creep deformation at 1423 K are shown in Figs. 12
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and 13. Figure 12 shows the strong dislocationparticle interaction in PM~TZM after creep tests at 200 MPa (Fig. 12(a)) and 375 MPa (Fig. 12(b)). Most dislocations are pinned by particles, which may be expected from the relatively low stress used. The carbides are smaller than those observed in the asreceived state (Fig. 2(a)). Therefore, these particles were probably formed during the creep tests. Annealing of the as-received specimens also caused a fine dispersion of particles (4-6 nm in diameter). Thus these new particles have their origin in the residual supersaturation of C at the testing temperature. The newly created particles are somewhat larger after creep at higher stresses. As most particles are found close to or at dislocation arrangements (walls, subgrain boundaries, etc.), strain-aided precipitate coarsening is likely to occur. In all cases, a very strong particle-dislocation interaction is seen (see arrows). In Fig. 12(b) some of the dislocations in the subgrain boundary are not visible under the diffraction conditions used, and particles are seen more clearly. Figure 13 shows particle-dislocation interactions in VM-TZM after creep deformation at 250 (Fig. 13(a)) and 350 MPa (Fig. 13(b)). At the lower stress, the
Figure 14 shows the fatigue response of PM-TZM in two microstructural conditions (after 70% and 90% cold work) and of VM-TZM (after 90% cold work) for tests performed at room temperature and at 1123 K. In the figure, N indicates the number of cycles necessary to fracture a specimen cyclically loaded to a stress
Fig. 12. Particle-dislocationinteraction in PM-TZMafter creep deformation:(a) creep stress 200 MPa; (b) creep stress 375 MPa.
Fig. 13. Particle-dislocationinteraction in VM-TZM alloy after creep deformation: (a) creep stress 250 MPa; (b) creep stress 350 MPa.
incoherent particles are also very effective for pinning dislocations. No dislocation loops are observed which would be required if the Orowan mechanism [11] had been active. Tilting experiments suggest that the dislocation-particle interaction results in attachment of the dislocation to the particle. This type of interaction resembles the mechanism proposed by Srolovitz et al. [12] for oxide dispersion strengthened alloys. In Fig. 13(a) at least five dislocations are in contact with a particle. This number is variable but dislocations are always effectively pinned. Figure 13(b) shows another case where several particles interact with dislocations in a subgrain boundary. Again, no loops are visible although the dislocations are bent almost 180" (see arrows in figures). 3.3. Fatigue experiments
H. A. CaMeron et al.
_ o a. The effect of prior cold work is readily apparent in the test performed at room temperature; an increase in fatigue life is observed with a higher amount of cold work. In addition, Fig. 14 shows that for the same degree of cold work, the fatigue limit at room temperature is similar for the two TZM alloys. This limit lies at approximately 500 MPa. For stress amplitudes higher than the endurance limit, the number of cycles to fracture for VM-TZM is higher than for PM-TZM. Figure 14 also shows results of fatigue tests at 1123 K. At this temperature, no fatigue limit is observed for the stress amplitudes used, and little difference can be found between the results for the two TZM alloys. These results are also independent of the amount of prior cold work, as shown in the same figure, and of the orientation of the specimens with respect to the prior deformation direction. Figure 15 shows the results of fatigue experiments obtained with welded specimens at room temperature and at 1123 K. These samples were electron beam welded. As in the base materials, a fatigue limit is observed at room temperature. However, in this case the endurance limit of VM-TZM is clearly higher (approximately 380 MPa) than that of PM-TZM (around 250 MPa). At 1123 K, no fatigue limit is observed for the stress amplitudes used (lower amplitudes are irrelevant for the technological applications of this alloy) and there is virtually no difference between the results of the two TZM alloys. As a comparison between Figs. 14 and 15 shows, the endurance
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limit of the base material at room temperature is about 1.5-2 times higher than that of the welded material. As for the high temperature tests, the ratio of fracture stresses in the matrix and the weld, O'amatrix/O'aweld, is a function of the number of cycles to failure and varies between approximately 2.9 for high stress amplitudes (N= 104) and 2.5 (N= 10 7) for lower amplitudes.
4. Discussion The two industrially applied processes to produce the Mo-base TZM alloy yield two basically different materials. Differences in grain and subgrain size have been found as well as differences in particle chemistry and size. The PM-TZM alloy contains large oxide particles which could act as sites for crack nucleation, owing to their brittleness and poor bonding to the matrix. The VM-TZM alloy contains only incoherent ZrxTil_xC particles, but they are considerably larger than the coherent TiC found in the PM-TZM alloy. The creep results can be rationalized on the basis of the microstructural information. The creep response in the center of the 50 mm VM-TZM bar is better than in the peripheral region. Here the effect of prior cold work and of particle size can be clearly observed. Large subgrain size, low dislocation density and large particles cause a decrease in creep strength for this section of the as-received bar. The creep properties of PM-TZM are more homogeneous throughout the 50 mm bar, as may be expected from the microstruc-
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Two Mo-base alloys
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ture. (Agronov et al. [7] also found a strong influence of prior cold work on the creep properties of PM-TZM.) However, the creep properties of PM-TZM and the center of the VM-TZM bar are very similar (see Fig. 11). Apparently, the combination of a fine distribution of carbides (those present at the beginning of the test and those formed during creep deformation) and a high dislocation density produces a very stable alloy under compressive creep conditions. Little effect of the large Zr20 particles was observed, but this can be related to the testing conditions. In compression, the brittle and debonded Zr20 particles can only impair the creep response marginally since crack growth is considerably reduced. In VM-TZM, no new particles are found after creep. The existing particles do not coarsen very much during a test of 150 h. Although the particles are larger than in PM-TZM, the creep response of the central region of the VM-TZM bar is comparable and even slightly better than that of the central region of the PMTZM bar. The interaction between particles and dislocations is very strong. The dislocations are effectively pinned by both sets of particles, i.e. the coherent TiC and the incoherent ZrxTil_xC particles. Two mechanisms may be evoked to account for the particle-dislocation interaction, the Orowan mechanism [11] and the mechanism proposed by Srolovitz et aL [12]. In the latter mechanism, dislocations are attracted to the particle-matrix interface and are forced to remain there until a given stress level is reached and the dislocations are freed. As Figs. 12 and 13 show, many dislocations seem to be attached to the particle-matrix interface. However, no dislocation loops around particles are observed. Therefore the present observations show that the Orowan mechanism cannot be active and rate controlling since dislocation loops necessarily result as a consequence of dislocation bending around the particles. Nevertheless, the stresses required to overcome a barrier are very similar for the two mechanisms. The stress limit for the Orowan mechanism lies between about 205 and 230 MPa for VMTZM and ranges from about 215 to 240 MPa for PM-TZM (for particle spacings of (490 + 20) /k and (470 _+20) A respectively). These values are obtained using the expressions given by Bacon et al. [13] and Lund and Nix [14]. As to the mechanism proposed by Srolovitz et al. [12], the critical stress to pull a dislocation away from a particle is 174 MPa for VM-TZM and 205 MPa for PM-TZM. These values are closer to the experimental stress value (180 MPa) below which the creep rate could not be measured. However, the possibility of interference of other mechanisms cannot be discarded. In particular, the effect of grain boundary sliding could be important at relatively low applied stresses.
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Two Mo-base alloys
The fatigue performance of the two T Z M alloys is similar. However, after electron beam welding, the fatigue properties of PM-TZM are affected more drastically. This is clearly related to the microstructure, as a large number of carbides and recrystallization produce a weak weld. In addition the presence of ZrO 2 can affect the weld strength, particularly during the tensile part of cyclic deformation.
5. Concluding remarks The TZM alloys produced commercially are diffegent in grain and subgrain size, precipitate size and chemical composition. The creep data obtained at 1423 K can be fitted to a power law. Stress exponents of 11, 15 and 14 are found for the periphery and center of the PM-TZM alloy bar and for the center of the VM-TZM alloy bar. The periphery of the VMTZM alloy bar showed poor creep properties. These differences can be explained on the basis of the observed microstructure. TEM observations suggest that the capture of dislocations at the interfaces between precipitates and matrix plays an important role in controlling the creep deformation. The friction-welding procedure produces a negligible decrease in the mechanical properties of the base alloy. The fatigue behavior of the two TZM alloys is similar but the welds of PM-TZM are weaker. Endurance limits of 380 MPa and 250 MPa were measured for the welds of VM-TZM and PM-TZM respectively.
Acknowledgments The authors are grateful to Mrs. R. B~inninger and Mr. A. Blanchard for experimental assistance.
References 1 W. Jakobeit, High Temp. High Pressures, 14(1982) 111. 2 G. Raule and R. Bauer, Proc. High temperature metallic materials for applications in gas cooled reactors, OEFZS. Rep. 4086, 1981 (OEFZS, Vienna), p. 01. 3 W. Jakobeit, J. P. Pfeiffer and G. Ullrich, Nucl. TechnoL, 66 (1984) 195. 4 R. Eck, Planseeber. Pulvermetall., 27(1979) 53. 5 W. Jacobeit, W. Bulla, R. Eck and G. Ullrich, Proc. 11th Int. PlanseeSeminar '85, Vol. 2, Reutte, Austria, 1985, p. 17. 6 W. Jacobeit, Int. J. Refract. HardMet., 2(1983) 133. 7 D. Agronov, E. Freund and A. Rosen, High Temp. High Pressures, 18(1986) 151. 8 F. Matsuda, M. Ushio, K. Nakata and Y. Edo, Trans. Jpn. Welding R es. Inst., 8 (1979) 57. 9 J. Wadsworth, G. R. Morse and P. M. Chewey, Mater Sci. Eng., 59(1983) 257-273.
H. A. Calderon et al. 10 W. P. Pearson, Handbook of Lattice Spacings and Structure of Metals andAlloys, Vol. 1, Pergamon, Oxford, 1958, p. 945. 11 E. Orowan, Symp. on Internal Stresses in Metals and Alloys, Session III Discussion, Institute of Metals, London, 1948, p. 451.
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12 D. J. Srolovitz, R. A. Petrovitz-Luton and M. J. Luton, Acta Metall., 31 (1983) 2151. 13 D. J. Bacon, U. E Kocks and R. O. Scattergood, Philos. Mag., 28(1973) 1241. 14 R.W. Lund and W. D. Nix, Acta Metall., 24(1976) 469.
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Microstructure and plasticity of two molybdenum-base alloys (TZM) H. A. Calderon and G. Kostorz Institut fiir Angewandte Physik, E TH Ziirich, CH-8093 Ziirich (Switzerland)
G. Ullrich Paul-Scherrer-Institut, CH-5232 Villigen PSI (Switzerland) (ReceivedMay 10, 1991; in revisedform September22, 1992)
Abstract In two different commercial Mo-base alloys (TZM), produced by vacuum melting and by powder metallurgy respectively, microstructural differences (particle size and chemistry, grain and subgrain sizes, dislocation density) have been found which affect the observed mechanical properties of the material. The compression-creep properties at 1423 K show a negligibly small creep rate at a stress of approximately 200 MPa. Trapping of dislocations by particles is proposed to be the controlling deformation mechanism during creep. The microstructure and fatigue properties of TZM welds were also investigated. Friction welds showed the best mechanical properties. Fatigue measurements in load control at room temperature and 1123 K show that the endurance limit of the vacuum-melted alloy is higher than that of the powdermetallurgically processed alloy.
1. Introduction Mo-base alloys are commonly used for applications at high temperatures where a high mechanical strength is required. The creep strength of Mo-base alloys at temperatures higher than l l 0 0 K is considerably better than that of Ni alloys [1]. Nevertheless, the poor oxidation resistance at temperatures higher than 750 K limits the use of these alloys to well controlled environments. Among the Mo alloys of technological interest, the T Z M alloy is the most important. Its applications range from turbine blades to high-temperature forging dies. The known good mechanical properties of TZM also made it attractive for application in a project involving high temperature helium turbines [2, 3]. The name of this alloy is related to its composition; Mo-0.5Ti-0.08Zr-(0.1-0.4)C (in atomic per cent). The presence of C leads to the formation of a hardening dispersion of carbides which strengthens the alloy beyond the limits shown by single-phase Mo alloys, i.e. solid-solution hardened Mo alloys containing W or Re [4]. T Z M is commercially produced via two industrial processes, vacuum melting (VM-TZM) and powder metallurgy (PM-TZM). The latter fabrication method is considerably less expensive but some reports [5, 6] indicate that the creep-rupture properties of PM-TZM are inferior to those of VM-TZM. In extensive investigations of the creep properties of PM-TZM conducted by Rosen and coworkers [7], however, the effects of 0921-5093/93/$6.00
particle dispersion and thermomechanical treatment on the creep properties were studied, and an excellent creep strength was found. It was concluded that the creep strength stems from a preferential nucleation of precipitates at subgrain boundaries. However, most of the details of the relevant hardening mechanisms and of the microstructure of this alloy are still unknown. In the present work, some creep properties of the two different T Z M alloys were studied. Furthermore, the relevant microstructures of both alloys were compared in the as-received state and after creep. From these observations the mechanism likely to control the creep rate may be identified. Several investigations were also conducted in order to study the conditions under which T Z M can be welded, with suitable mechanical properties of the welds (see for example refs. 8 and 9). In general, most industrial processes produce welds with inferior mechanical properties. This is due to recrystallization and pore formation induced by the welding process [8, 9]. Some fatigue properties of welded and base materials are reported. Finally, the microhardness and the microstructures produced by two welding procedures are compared.
2. Experimental procedures The specimens for the study of the creep properties in uniaxial compression were obtained from bars © 1993 - Elsevier Sequoia. All fights reserved
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50 mm in diameter by cutting pieces of a square crosssection with an aspect ratio lid of approximately 2 (here I is the length of the specimen and d the equivalent diameter), in order to minimize buckling and barreling. A thin layer of BN was used as a lubricant. All tests were performed in a screw-driven machine programmed to keep the applied stress constant. The testing temperature was 1423 K, and a vacuum better than 5 x 10-3 Pa was maintained during the tests. Transmission electron microscopy (TEM) was performed in a JEOL 100C microscope. Thin foils were produced by electropolishing in a twin-jet apparatus with a solution of 12.5 vol.% H2SO 4 in ethanol at 258 K. A voltage of 10 V produced a current density of approximately 3 mA mm- 2. Microstructures after welding were observed after either electron beam welding two plates of a thickness of 8 mm or friction welding bars of a diameter of 25 mm. The specimens were stress relieved at 973 K f o r l h. The fatigue tests were performed in a screw-driven machine (Schenk, Type PH 1G 0016) in load control at a frequency of 25 Hz (sinusoidal waveform), with a ratio O'max/Ormi n = - 1. The fatigue specimens had a cylindrical gage section with a diameter of 1.7 mm or 2 mm. The gage length was five times the diameter. The high-temperature (T = 1123 K) tests took place in a helium atmosphere that prevented the production of volatile Mo oxides. This atmosphere simulated the environment of a high-temperature gas-cooled reactor. It was also possible to add controlled amounts of impurities to the helium atmosphere.
Two Mo-base alloys TABLE 1. Grain and subgrain sizes in TZM alloys Location
PM-TZM l (/~m) (/~m)
Grain size Edge of bar 38 Center of bar 33 Subgrain size Edge of bar Center of bar
1.2 1
114 128
VM-TZM R=lgl¢v ¢v l (#m) (/~m)
R=~I~
3 3.9
14.9 15.6
0.7 1.7 0.74 1.4
19 26 3.6 5.1
284 406 0.9 0.8
4 6.4
3. Experimental results 3.1. Microstructural observations 3.1.1. As-received specimens The two TZM alloys have different grain structures, as summarized in Table 1. The grains of PM-TZM are elongated parallel to the extrusion direction. The aspect ratio ]g/~ (/g and ~ represent the mean length and thickness respectively) is relatively low. The aspect ratio for VM-TZM is considerably higher than for PMTZM. In PM-TZM, coarse particles can be seen as shown in Fig. l(a), their sizes ranging from 1 to 15/am. Some of these particles are cracked and exhibit poor bonding to the matrix. Energy-dispersive X-ray fluorescence analysis shows that these large particles contain mostly Zr with minor amounts of Ti and Ca. Electron diffraction reveals a cubic structure of the particles, with a lattice parameter of around 5.17 A. Therefore, it is very likely that the particles are ZrO2 which could be stabilized at room temperature by Ca. The grain struc-
Fig. 1. Grain microstructuresof the as-received samples, parallel to the extrusion direction: (a) PM-TZM (the dark, coarse particles are ZrO2); (b) VM-TZM.
ture of VM-TZM is typical of cold-worked materials (Fig. l(b)). No coarse particles are found in this case. TEM also reveals considerable differences in the microstructure of the two as-received TZM alloys. The subgrain size and aspect ratio, the dislocation arrangement and the carbide distribution and size are all characteristic of each alloy. Figure 2 shows typical particle populations in the two alloys. The particles in PM-TZM (Fig. 2(a)) are in general homogeneously distributed in subgrain boundaries and in the subgrain
H. A. CaMeron et al.
interiors. The particle contrast suggests that precipitates are coherent or semicoherent and spherical. Considering the alloy composition, and the fact that most of the Zr forms ZrO 2 (Fig. l(a)), the small particles in Fig. 2(a) should be TiC. This is supported by TEM observations of the strain contrast expected for such particles. However, the average particle sizes are not independent of the specimen location in the 50 mm bar. A summary of the measurements is given in Table 2. In VM-TZM, precipitate size and morphology are considerably different from those of PM-TZM (Fig. 2(b)). The particles are not perfectly spherical and are apparently incoherent or only partially coherent with the matrix. Strain contrast is observed only with highorder reflections. The dislocation structure in this alloy shows a higher degree of recovery. Some dislocation networks can be observed in an advanced degree of development. Furthermore, many particles are connected by dislocations which probably promotes accelerated coarsening of particles. Table 2 summarizes the metallographic measurements. It is important to note that particle sizes and dislocation densities
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are not constant across the diameter of the original bar, which can be linked to the observed variation in mechanical behavior (see below). In VM-TZM, no large oxide particles are found and therefore the particle chemistry should be ZrxTi,_xC. The lattice mismatch complete coherency, thus reducing between Mo and ZrxTi,_~C is large enough to prevent strain contrast in TEM for low-order reflections. The exact amount of Ti in these carbides was not determined experimentally. However, measurements of lattice parameters in TEM and comparison with published data [10] indicate that x is approximately 0.5-0.6. The volume fraction fv of precipitates can be calculated from the nominal composition to range from 0.24% to 0.36%. Experimentally, fv = 1%-1.5% was found. 3.1.2. Welds Microhardness measurements. Figures 3 and 4 show
the microhardness profiles through the welded areas in the two alloys after electron-beam and friction welding. All measurements were taken in a direction perpendicular to the weld. In both alloys the microhardness values of the welded regions and of the heat-affected zones remain constant or show little change after friction welding. In contrast, electron-beam welding causes
TABLE 2. Dislocation density and average particle sizes in TZM alloys Location
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VM-TZM
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Fig. 3. Microhardness (HV) profiles taken across the welds of PM-TZM: x electron-beam weld; • friction weld.
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a decrease in microhardness by about 30% in both TZM alloys. The quasi-homogeneous microhardness observed for the friction-welded alloys indicates the relatively low damage induced by this welding technique, whereas a considerable reduction in strength is brought about by electron-beam welding. The absolute
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Two Mo-base alloys
values of microhardness depend on the thermomechanical history of the specimen and will not be discussed here.
Microstructure after electron-beam welding. The grain structure of the weld region in both TZM alloys shows large differences in grain size as shown in Fig. 5(a)-(c) for VM-TZM. In this figure, the weld (Fig. 5(a)), the heat-affected zone (HAZ) (Fig. 5(b)-(c)) and the base material (Fig. 5(c)) can be observed. Recrystallization in the weld and the H A Z is evident. Similar features can be seen in the case of PM-TZM but, in addition, there are large pores in the weld, with sizes ranging from 20/~m to 200/~m (Fig. 5(d)). The size of the H A Z is approximately 12 mm in both cases. Furthermore, transgranular cracks were found in the VM-TZM weld. This is probably due to an inappropriate welding atmosphere [8]. Table 3 shows the grain sizes measured in the weld and the HAZ. Figure 6 shows the grain microstructure as observed in the scanning electron microscope. The weld of PMTZM has large particles and pores inside the grains. The bright areas that decorate the grain and subgrain boundaries are most probably not an artefact. Subse-
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Fig. 5. Grain structure after electron-beam welding:(a) VM-TZM weld; (b) VM-TZMheat affected zone; (c) VM-TZMbase material; (d) PM-TZMweld.
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TABLE 3. Grain sizes in welds and heat affected zones (HAZ) Alloy
Processa
Region
Grain size (mm)
PM-TZM
EBW
Weld
PM-TZM VM-TZM
EBW EBW
HAZ Weld
VM-TZM PM-TZM PM-TZM VM-TZM VM-TZM
EBW FW FW FW FW
HAZ Weld HAZ Weld HAZ
0.25 width 0.80 length 0.70 0.25 width 0.75 length 0.060 0.008 0.025 0.007 0.030
aEBW electron beam welding, FW friction welding.
Fig. 6. Grain structure of PM-TZM after an electron-beam welding as observed by SEM. The particles at the grain and subgrain boundaries are presumably (Ti,Mo)2C. The arrowed particles are ZrO 2. quent TEM studies revealed a high particle density at grain boundaries. These particles are probably Mo carbides (see below). The larger particles marked by arrows are Zr20 as described previously. TEM shows that the dislocation density in the weld and the H A Z of PM-TZM is very low, typical of a recrystallized structure. The dislocation density is relatively high only around large particles. Very few semicoherent TiC particles are observed in the weld. Instead, large carbide particles rich in Ti and Mo are visible along the grain and subgrain boundaries. These particles have sizes ranging from 250 nm to 2/tm. Wadsworth et al. [9] have identified similar particles as (Ti,Mo)2C. The weld in VM-TZM also shows a low dislocation density (less than 107 cm-2). Only large, elongated (Ti,Mo)2C particles are observed, mostly along grain boundaries. The finer ZrxTil_xC particles found in the as-received alloy are not present. Figure 7(a) illustrates the grain boundary carbides, and Fig.
Fig. 7. (a) Grain boundary carbides in the weld region of VMTZM after friction welding; (b) corresponding diffraction pattern. 7(b) shows the corresponding diffraction pattern. The H A Z of the VM-TZM weld shows elongated subgrains, thus recrystallization did not occur in this case. Areas with large ZrxTil_xC particles are found. As in the as-received alloy, strong dislocation-particle inte"actions are observed. The large (Ti, Mo)2C particles are also observed along the grain boundaries, but less frequently than in the PM-TZM alloy. Microstructure after friction welding. The weld has a fine-grained structure in both alloys as shown in Fig. 8. The extension of the welds and H A Z s is smaller than 1.5 mm. This is considerably less than for alloys which are electron beam welded. In the PM-TZM weld, a
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relatively high number of particles (size 10-15/am) can be seen (Fig. 8(a)). These particles are rich in Ti and/or Mo, and most of them are carbides. Table 3 lists the measurements of grain size. TEM shows that the H A Z of PM-TZM has elongated subgrains (2-2.5 /am in length) with large Tiand/or Mo-rich carbides. Some signs of recrystallization can also be seen but this process is not complete, as indicated by frequently occurring areas of high dislocation density. Figure 9 shows representative views of the subgrain microstructure and the carbides in the weld. It can be observed (Fig. 9(b)) that hardening carbides (TIC) are still present within the welded region, increasing the mechanical strength of the weld in comparison with electron-beam welding. The subgrains in the VM-TZM weld are equiaxed with an average subgrain size of 2/am. In the weld, recrystallization is incipient, and areas with a high dislocation density can be observed. A TEM determination yields a dislocation density of 7 x 10 ~° cm -2 for these areas. This dislocation density is higher than that found in PM-TZM. Large carbides (1-2/am) are also, but infrequently, observed along subgrain boundaries. A few incoherent ZrxTil_xC particles are observed inter-
Fig. 8. Grain structure after friction welding: (a) PM-TZM (the large particles marked by arrows are M2C-typecarbides rich in Mo and Ti); (b) VM-TZM.
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Two Mo-base alloys
acting with and pinning dislocations. In the H A Z the subgrains vary more broadly in size, and some of them do not seem to be equiaxed. Very few large carbides are found. Instead, a fine homogeneous distribution of ZrxTi~_xC is observed. The mean size of these particles is 50 nm. These observations suggest that the grains and subgrains produced by the welding procedure are very fine and equiaxed and that no recrystallization takes place in these regions. Some deformation is also detectable. In both welds, large oxide particles are present. Analysis of diffraction patterns shows that they correspond to monoclinic MOO2. A dark-field view of these oxides in PM-TZM and the corresponding diffraction pattern with a zone axis close to [110] are shown in Fig. 10. The objective aperture used for the dark-field image (Fig. 10(b) was large enough to produce phase contrast with the transmitted beam and as MoO2 has large lattice constants, lattice fringes are visible for the (100) plane. In the HAZ, the incipient recrystallization produces a variable subgrain size and also precipitation of large carbide particles which are clearly deleterious for the
Fig. 9. Weld region of PM-TZM after friction welding: (a) subgrains and carbides;(b) TiC particles.
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mechanical properties of the weld. Hardening carbides (i.e. TiC for PM-TZM and TixZr ]_ xC in VM-TZM) are observed more frequently than in the weld.
3.2. Creep experiments 3.2.1. Creep rate Figures l l(a) and l l(b) show the stationary creep strain rate g at 1423 K as a function of the constant compressive stress used experimentally for the two alloys tested. In the figures, the positions of the specimens in the original bar are indicated by numbers, 5 representing the center and 1 the bar edge. Results for specimens 1 are not reported as the data scattered considerably. Figure 1 l(a) shows that there is a variation in the steady-state creep rate with specimen location for the VM-TZM bar. The peripheral region
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Fig. 11. Steady-state creep rate g vs. applied stress a in TZM alloys at 1423 K: (a) VM-TZM; (b) PM-TZM. The numbers indicate the specimen location in the original bar of a diameter of 50 mm; 1 indicates the bar edge and 5 the center.
shows a lower creep strength. In PM-TZM, the creep response is more uniform, and only small differences in g can be found. The data can be fitted to a power law oc a" with good data correlation. The stress exponent n takes the values 15 and 11 for the center and periphery of the PM-TZM bar and 14 for the center of the VM-TZM bar. At high stresses, the creep response of the central region of VM-TZM appears to be slightly better than that of the center of the PM-TZM bar. The steady-state creep rate induced by stresses lower than 180 MPa was below the resolution limit of the extensometer used.
3.2.2. Microstructure after creep deformation Representative views of the microstructure after creep deformation at 1423 K are shown in Figs. 12
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Two Mo-base alloys
and 13. Figure 12 shows the strong dislocationparticle interaction in PM~TZM after creep tests at 200 MPa (Fig. 12(a)) and 375 MPa (Fig. 12(b)). Most dislocations are pinned by particles, which may be expected from the relatively low stress used. The carbides are smaller than those observed in the asreceived state (Fig. 2(a)). Therefore, these particles were probably formed during the creep tests. Annealing of the as-received specimens also caused a fine dispersion of particles (4-6 nm in diameter). Thus these new particles have their origin in the residual supersaturation of C at the testing temperature. The newly created particles are somewhat larger after creep at higher stresses. As most particles are found close to or at dislocation arrangements (walls, subgrain boundaries, etc.), strain-aided precipitate coarsening is likely to occur. In all cases, a very strong particle-dislocation interaction is seen (see arrows). In Fig. 12(b) some of the dislocations in the subgrain boundary are not visible under the diffraction conditions used, and particles are seen more clearly. Figure 13 shows particle-dislocation interactions in VM-TZM after creep deformation at 250 (Fig. 13(a)) and 350 MPa (Fig. 13(b)). At the lower stress, the
Figure 14 shows the fatigue response of PM-TZM in two microstructural conditions (after 70% and 90% cold work) and of VM-TZM (after 90% cold work) for tests performed at room temperature and at 1123 K. In the figure, N indicates the number of cycles necessary to fracture a specimen cyclically loaded to a stress
Fig. 12. Particle-dislocationinteraction in PM-TZMafter creep deformation:(a) creep stress 200 MPa; (b) creep stress 375 MPa.
Fig. 13. Particle-dislocationinteraction in VM-TZM alloy after creep deformation: (a) creep stress 250 MPa; (b) creep stress 350 MPa.
incoherent particles are also very effective for pinning dislocations. No dislocation loops are observed which would be required if the Orowan mechanism [11] had been active. Tilting experiments suggest that the dislocation-particle interaction results in attachment of the dislocation to the particle. This type of interaction resembles the mechanism proposed by Srolovitz et al. [12] for oxide dispersion strengthened alloys. In Fig. 13(a) at least five dislocations are in contact with a particle. This number is variable but dislocations are always effectively pinned. Figure 13(b) shows another case where several particles interact with dislocations in a subgrain boundary. Again, no loops are visible although the dislocations are bent almost 180" (see arrows in figures). 3.3. Fatigue experiments
H. A. CaMeron et al.
_ o a. The effect of prior cold work is readily apparent in the test performed at room temperature; an increase in fatigue life is observed with a higher amount of cold work. In addition, Fig. 14 shows that for the same degree of cold work, the fatigue limit at room temperature is similar for the two TZM alloys. This limit lies at approximately 500 MPa. For stress amplitudes higher than the endurance limit, the number of cycles to fracture for VM-TZM is higher than for PM-TZM. Figure 14 also shows results of fatigue tests at 1123 K. At this temperature, no fatigue limit is observed for the stress amplitudes used, and little difference can be found between the results for the two TZM alloys. These results are also independent of the amount of prior cold work, as shown in the same figure, and of the orientation of the specimens with respect to the prior deformation direction. Figure 15 shows the results of fatigue experiments obtained with welded specimens at room temperature and at 1123 K. These samples were electron beam welded. As in the base materials, a fatigue limit is observed at room temperature. However, in this case the endurance limit of VM-TZM is clearly higher (approximately 380 MPa) than that of PM-TZM (around 250 MPa). At 1123 K, no fatigue limit is observed for the stress amplitudes used (lower amplitudes are irrelevant for the technological applications of this alloy) and there is virtually no difference between the results of the two TZM alloys. As a comparison between Figs. 14 and 15 shows, the endurance
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limit of the base material at room temperature is about 1.5-2 times higher than that of the welded material. As for the high temperature tests, the ratio of fracture stresses in the matrix and the weld, O'amatrix/O'aweld, is a function of the number of cycles to failure and varies between approximately 2.9 for high stress amplitudes (N= 104) and 2.5 (N= 10 7) for lower amplitudes.
4. Discussion The two industrially applied processes to produce the Mo-base TZM alloy yield two basically different materials. Differences in grain and subgrain size have been found as well as differences in particle chemistry and size. The PM-TZM alloy contains large oxide particles which could act as sites for crack nucleation, owing to their brittleness and poor bonding to the matrix. The VM-TZM alloy contains only incoherent ZrxTil_xC particles, but they are considerably larger than the coherent TiC found in the PM-TZM alloy. The creep results can be rationalized on the basis of the microstructural information. The creep response in the center of the 50 mm VM-TZM bar is better than in the peripheral region. Here the effect of prior cold work and of particle size can be clearly observed. Large subgrain size, low dislocation density and large particles cause a decrease in creep strength for this section of the as-received bar. The creep properties of PM-TZM are more homogeneous throughout the 50 mm bar, as may be expected from the microstruc-
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ture. (Agronov et al. [7] also found a strong influence of prior cold work on the creep properties of PM-TZM.) However, the creep properties of PM-TZM and the center of the VM-TZM bar are very similar (see Fig. 11). Apparently, the combination of a fine distribution of carbides (those present at the beginning of the test and those formed during creep deformation) and a high dislocation density produces a very stable alloy under compressive creep conditions. Little effect of the large Zr20 particles was observed, but this can be related to the testing conditions. In compression, the brittle and debonded Zr20 particles can only impair the creep response marginally since crack growth is considerably reduced. In VM-TZM, no new particles are found after creep. The existing particles do not coarsen very much during a test of 150 h. Although the particles are larger than in PM-TZM, the creep response of the central region of the VM-TZM bar is comparable and even slightly better than that of the central region of the PMTZM bar. The interaction between particles and dislocations is very strong. The dislocations are effectively pinned by both sets of particles, i.e. the coherent TiC and the incoherent ZrxTil_xC particles. Two mechanisms may be evoked to account for the particle-dislocation interaction, the Orowan mechanism [11] and the mechanism proposed by Srolovitz et aL [12]. In the latter mechanism, dislocations are attracted to the particle-matrix interface and are forced to remain there until a given stress level is reached and the dislocations are freed. As Figs. 12 and 13 show, many dislocations seem to be attached to the particle-matrix interface. However, no dislocation loops around particles are observed. Therefore the present observations show that the Orowan mechanism cannot be active and rate controlling since dislocation loops necessarily result as a consequence of dislocation bending around the particles. Nevertheless, the stresses required to overcome a barrier are very similar for the two mechanisms. The stress limit for the Orowan mechanism lies between about 205 and 230 MPa for VMTZM and ranges from about 215 to 240 MPa for PM-TZM (for particle spacings of (490 + 20) /k and (470 _+20) A respectively). These values are obtained using the expressions given by Bacon et al. [13] and Lund and Nix [14]. As to the mechanism proposed by Srolovitz et al. [12], the critical stress to pull a dislocation away from a particle is 174 MPa for VM-TZM and 205 MPa for PM-TZM. These values are closer to the experimental stress value (180 MPa) below which the creep rate could not be measured. However, the possibility of interference of other mechanisms cannot be discarded. In particular, the effect of grain boundary sliding could be important at relatively low applied stresses.
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The fatigue performance of the two T Z M alloys is similar. However, after electron beam welding, the fatigue properties of PM-TZM are affected more drastically. This is clearly related to the microstructure, as a large number of carbides and recrystallization produce a weak weld. In addition the presence of ZrO 2 can affect the weld strength, particularly during the tensile part of cyclic deformation.
5. Concluding remarks The TZM alloys produced commercially are diffegent in grain and subgrain size, precipitate size and chemical composition. The creep data obtained at 1423 K can be fitted to a power law. Stress exponents of 11, 15 and 14 are found for the periphery and center of the PM-TZM alloy bar and for the center of the VM-TZM alloy bar. The periphery of the VMTZM alloy bar showed poor creep properties. These differences can be explained on the basis of the observed microstructure. TEM observations suggest that the capture of dislocations at the interfaces between precipitates and matrix plays an important role in controlling the creep deformation. The friction-welding procedure produces a negligible decrease in the mechanical properties of the base alloy. The fatigue behavior of the two TZM alloys is similar but the welds of PM-TZM are weaker. Endurance limits of 380 MPa and 250 MPa were measured for the welds of VM-TZM and PM-TZM respectively.
Acknowledgments The authors are grateful to Mrs. R. B~inninger and Mr. A. Blanchard for experimental assistance.
References 1 W. Jakobeit, High Temp. High Pressures, 14(1982) 111. 2 G. Raule and R. Bauer, Proc. High temperature metallic materials for applications in gas cooled reactors, OEFZS. Rep. 4086, 1981 (OEFZS, Vienna), p. 01. 3 W. Jakobeit, J. P. Pfeiffer and G. Ullrich, Nucl. TechnoL, 66 (1984) 195. 4 R. Eck, Planseeber. Pulvermetall., 27(1979) 53. 5 W. Jacobeit, W. Bulla, R. Eck and G. Ullrich, Proc. 11th Int. PlanseeSeminar '85, Vol. 2, Reutte, Austria, 1985, p. 17. 6 W. Jacobeit, Int. J. Refract. HardMet., 2(1983) 133. 7 D. Agronov, E. Freund and A. Rosen, High Temp. High Pressures, 18(1986) 151. 8 F. Matsuda, M. Ushio, K. Nakata and Y. Edo, Trans. Jpn. Welding R es. Inst., 8 (1979) 57. 9 J. Wadsworth, G. R. Morse and P. M. Chewey, Mater Sci. Eng., 59(1983) 257-273.
H. A. Calderon et al. 10 W. P. Pearson, Handbook of Lattice Spacings and Structure of Metals andAlloys, Vol. 1, Pergamon, Oxford, 1958, p. 945. 11 E. Orowan, Symp. on Internal Stresses in Metals and Alloys, Session III Discussion, Institute of Metals, London, 1948, p. 451.
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12 D. J. Srolovitz, R. A. Petrovitz-Luton and M. J. Luton, Acta Metall., 31 (1983) 2151. 13 D. J. Bacon, U. E Kocks and R. O. Scattergood, Philos. Mag., 28(1973) 1241. 14 R.W. Lund and W. D. Nix, Acta Metall., 24(1976) 469.