Dislocation structure in precipitation-hardened Ir-based binary alloys Y. Yamabe-Mitarai, Y. Ro, S. Nakazawa, T. Maruko,* and H. Harada National Research Institute for Metals, 1-2-1 Sengen, Tsukuba, Ibaraki, 305-0047, Japan *Furuya Metals Co. Ltd., 1915 Morizoeshima, Shimodate, Ibaraki, 308-0861, Japan Keywords: compressive strength, dislocation structure, shearing mode, bypass mode, cuboidal precipitates, maze structure, orderly arrayed structure ABSTRACT This paper examines the microstructure and the strength behavior up to 1800˚C of Ir-V, -Ti, -Nb, -Ta, -Hf, and -Zr alloys which formed fcc and L12 two-phase coherent structures. It is indicated that the precipitation-hardening effect depends on microstructure, resulting from differences in the lattice misfit. A maze structure such as in the Ir-Zr alloy was more effective for precipitation hardening than an orderly arrayed structure such as in the Ir-Nb alloy. The dislocation structure of deformed samples was investigated by TEM. The deformation mechanism of the alloys with different microstructures was discussed. 1. INTRODUCTION As high-temperature materials, platinum-group metals and their alloys have been investigated because of their high melting temperature and strength at high temperatures. Since Fleischer et al. suggested the potential of platinum-group metals-based intermetallic compounds as high-temperature structural materials [1, 2], several intermetallic compounds have been investigated. RuAl [3], RuSc [2], RuNb [2, 4], and IrAl [4] with a B2 structure, IrNb [1] and RuTa [2] with an L10 structure, and Ir3 Nb and Ir3 Zr [5] with an L12 structure have been noted. In addition to these intermetallic compounds, Ir-0.3wt%W alloys with an fcc single-phase microstructure have been noted as a cladding and post-impact containment material for radioactive fuel in radioisotope thermoelectric-generator heat sources [6]. Most of the research presented above concentrates on single-phase alloys. On the other hand, we have concentrated on fcc and L1 2 two-phase alloys because it is well known that L12 single-phase alloys, as well as fcc single-phase alloys, are less resistant to creep deformation than fcc and L12 two-phase alloys such as Ni-based superalloys [7, 8]. Thus, we have proposed a new class of superalloys, named refractory superalloys, based on Ir or Rh [9-12]. These superalloys have fcc and L12 two-phase coherent structures similar to those of Ni-based superalloys, but considerably higher melting temperatures. In our previous studies, the compression strengths of some of the as-cast Ir-based binary alloys were found to be over 800 MPa up to 1200˚C and about 200 MPa at 1800˚C [12]. We found an orderly arrayed structure in the Ir-Nb and Ir-Ta alloys after heat treatment. On the other hand, in the Ir-Hf and Ir-Zr alloys, a three dimensional maze structure formed. We indicated that differences in the microstructures depend on the lattice misfit between the fcc matrix and the L12 precipitates [11]. The strength behavior of Ir-Zr and Ir-Nb alloys has been investigated with respect to alloys with the maze structure and the orderly arrayed structure, respectively. It is indicated that the precipitation hardening of the Ir-Zr alloy is larger than that of the Ir-Nb alloy [13]. We observed shearing of precipitates in Ir-Zr alloys deformed at room temperature and 1200˚C [14]. Although the deformation structure of Ir-Nb alloys was not clear by observation with transmission electron microscopy (TEM), the bypass mechanism is suggested as the deformation mode [15]. In the present paper, an extended investigation was conducted for Ir-V, -Ti, -Ta, and -Hf binary alloys. All samples were heat-treated to produce L12 precipitates. A compression test up to 1200˚C was carried out on the heat-treated samples. The dislocation structures of deformed samples were observed by TEM. We compared these results with our previous results for the Ir-Nb and -Zr alloys and discussed the deformation mechanism. We also investigated the strength behavior and deformation mechanism at 1800˚C of Ir-based alloys, including Ir-Nb and Ir-Zr alloys, to understand their potential as high-temperature materials.
2. EXPERIMENTAL PROCEDURE Ir-X (X=V, Ti, Ta, Nb, Hf, or Zr) binary alloys were prepared as 70g button ingots by arc melting in an argon atmosphere. The alloy composition, component phases, and heat treatment are summarized in Table 1. Cylindrical samples with a diameter of 3 or 4 mm and a height of 6 or 8 mm were cut from the binary-alloy ingots for the compression tests and microstructural examination. Table 1 Alloy composition (at%), component phases, and heat treatment of tested alloys. Alloy composition Phase Heat treatment for compression test Ir-9Ti fcc 1200°C 1 hour Ir-15, 17Ti fcc +L12 1200°C 168 hours Ir-25Ti L12 2000°C 17 hours Ir-10Nb fcc 2000°C 72 hours Ir-15, 17Nb fcc +L12 1200°C 72, 168 hours, Ir-25Nb L12 2000°C 72 hours Ir-2Hf fcc 1800°C 1 hour Ir-13, 15Hf fcc +L12 1400 168 hours Ir-24Hf L12 2000°C 72 hours Ir-2Zr fcc 1800°C 1 hour Ir-12, 15Zr fcc +L12 1200°C 10 hours Ir-25Zr L12 2000°C 72 hours Ir-18Ta fcc +L12 1200°C 72 hours Ir-18V fcc +L12 1200°C 168 hours According to the phase diagrams [16], the application of solution heat treatment in the fcc single phase in Ir-based alloys is difficult, except for Ir-V alloys. Thus, to form precipitates, cylindrical as-cast samples were heated at temperatures between 1200 and 2000˚C in the fcc and L12 two-phase region for up to 168 hours. Heating at 1200˚C was carried out with these samples encapsulated in quartz tubes filled with argon gas, followed by water quenching. For heat treatments above 1200˚C, the whole sequence of heating and cooling was carried out within a vacuum furnace with a tungsten mesh heater to prevent damage of the heater by induction of air above 100 ˚C. The microstructures of the samples were observed by scanning electron microscopy (SEM), on a Philips model XL30, after electrolytically etching in an ethyl alcohol solution of 5% HCl. Thin discs with a thickness of 0.2 mm were cut from the heattreated samples, polished, and ion-milled for observation by transmission electron microscopy (TEM). The microstructure was observed using bright-field and dark-field imaging in a Philips model CM200 TEM operating at 200 kV. Compression tests were carried out on heat-treated samples in air for tests at temperatures up to 1200˚C and in an argon atmosphere for tests at 1800˚C. TENSILON/UTM-1-50000CW and INSTRON 8560 testing machines were used for compression tests up to 1200˚C and at 1800˚C, respectively. For high-temperature tests, the samples were kept at the testing temperature for 15 minutes before loading. The initial compressive strain rate for each test was 3.0 x 10-4/s. When the plastic strain reached 2-3%, the compression test was stopped in order to observe the dislocation structure in the deformed samples. The sample length was measured before and after the deformation to determine the real plastic strain. 3. RESULTS 3.1. PRECIPITATE MORPHOLOGY Typical dark-field or bright-field images of heat-treated Ir-based alloys are shown in Fig. 1. These phases were confirmed as the fcc and L12 structures by X-ray analysis [11] and selected area diffraction pattern in TEM. The dark-field images of the L12 phase were taken with a superlattice reflection (g=110). Thus, the phases with the L12 structure, namely, Ir3 V, Ir 3 Ti, Ir3 Nb, and Ir3 Zr, show bright contrast (Fig. 1a, b, c, and f, respectively). Bright-field images of the Ir-Ta and Ir-Hf alloys were taken (Fig. 1d and e) because detection of superlattice reflections from the L12 structures of these alloys is difficult due to the closer atomic-scattering
factors of Ir, Ta, and Hf, which result in very small structure factors of the superlattices in the L12 structures. The lattice misfit of the matrix and precipitates at 1200˚C, δ=( a p -a m)/a mx100, determined by high-temperature X-ray defractometry [11], are also shown in the photos as L.M. Here, the a p and a m are the lattice parameters of the precipitate and the matrix, respectively. When the lattice misfit was about –0.2 ~ 0.1% (Ir-V and -Ti alloys), irregular shaped precipitates were observed. The Ir-Ti alloy showed dislocation contrast in the precipitates (Fig. 1b). The misfit dislocation spacing was calculated using the lattice parameter of Ir3 Ti, (0.387 nm), and the lattice misfit at 1200˚C, (0.0012) [11], and found to be about 320 nm. The dislocation spacing observed was about 30 nm, which is one order of magnitude smaller than the calculated value. This shows that these dislocations are not misfit dislocations and the microstructure is therefore not a semi-coherent structure. We consider that these dislocations are formed in the precipitate and not at the interface. When the lattice misfit was about 0.3% (Ir-Nb alloy), cuboidal precipitates 100 nm in size, whose habit planes were {100}, formed and these cuboidal precipitates tended to align along the cube direction (Fig. 1c). We call this microstructure an orderly arrayed structure. Fine L1 2 precipitates 20 nm in size also formed in the fcc matrix during cooling in the Ir-Nb alloy. In the Ir-Ta alloy, precipitate shape and the existence of fine precipitates were not clear because of difficulty taking dark-field images. However, we expect a similar microstructure to the Ir-Nb alloy in the Ir-Ta alloy because these two alloys had similar lattice misfit. Precipitate size in the Ir-Ta alloy seems to be about 100 nm (Fig. 1d). When the lattice misfit was large, around 2% (Ir-Hf and Ir-Zr alloys), plate or rodlike precipitates 20 nm in width and from 100 to 300 nm in length were formed although we did not detect precipitate shape in the present investigation (Fig. 1e, f). The habit planes of the precipitates were {100}. These precipitates formed a three-dimensional maze structure in the IrZr alloy. How this maze structure formed is under investigation. We expect a similar microstructure to the Ir-Zr alloy in the Ir-Hf alloy, although we can not detect if the maze structure exists in the Ir-Hf alloy or not in this picture. If the precipitates are incoherent with the matrix, the precipitate shape is an oblate spheroid that balances interfacial energy and elastic strain energy [17]. We therefore consider cuboidal and plate or rod-like precipitates to be either perfectly coherent or semi-coherent because these precipitates are surrounded by {100} planes. Although the interface was observed using the tilting method in TEM, we did not observe misfit dislocations at the interface. Thus, we conclude that these precipitates are coherent. On the other hand, a different microstructure was also observed in the Ir-Hf and -Zr alloys with a large lattice misfit (Fig. 1g and h). In these alloys, precipitates with bright contrast lost their plate or rod-like shape. On the interface, we observed dislocations, which aligned at the interface with regular spacing.
Fig. 1 Precipitate shape of the Ir-15at% (a) V, (b) Ti, (c) Nb, (d) Ta, (e, g) Hf, and (f, h) Zr alloys. (a, b, c, f) are dark-field images. (d, e, g, h) are bright-field images. They were taken in the beam direction of [001]. Only (h) was taken in the beam direction of [101]. All dark-field images were taken with a superlattice reflection (g=110) from the L12 structure. Heat treatment at 1400˚C for 168 hours was carried out for the Ir-Hf alloy. Heat treatment at 1200˚C for 168 hours was carried out for other alloys. The misfit dislocation spacing of the Ir-Zr alloy was calculated to be about 18 nm using the lattice parameter of Ir 3 Zr, (0.395 nm), and the lattice misfit at 1200˚C, (0.022) [11]. The spacing of observed dislocations was about 13 nm. This shows that these dislocations are misfit dislocations. We conclude that this microstructure (Fig. 1g and h) is a semi-coherent structure. This shows that when the lattice misfit is large, both coherent and semi-coherent structures exist in the same sample. 3.2. STRENGTH BEHAVIOR
The temperature dependence of the compressive strength of the Ir-based alloys is shown in Fig. 2 with reference alloys [18-21]. The Ir-based alloys shown in this diagram have 50% L12 precipitates in volume fraction, and their compositions are determined as Ir18at% V, -17at% Ti, -17at% Nb, -18at% Ta, 13at% Hf, and -12at% Zr, respectively, using the phase diagram. The strengths of Ni-based superalloys (Ni-Al-Cr, MarM247, CMSX-10) decreased drastically with increasing the temperature above 800˚C. Although the strengths of the Ir-Nb, -Ta, -Hf, and -Zr alloys also decreased with increasing testing temperature, they were equivalent to or far higher than the strengths of MarM247 and CMSX-10 at room temperature and 1200˚C. Fig. 2 Temperature dependence of compressive On the other hand, the strengths of the Ir-Ti strength in Ir-based binary alloys with 50% and Ir-V alloys were low below 1200˚C. At precipitates in volume fraction and strengths of the 1800˚C, the strengths of the Ir-based alloys Ni- and W-based alloys. were between 100 and 200 MPa. This value is equivalent to that of W-based alloys, which was the highest at 1800˚C. Recently, Ma et al. showed that the strength of a Nb-Si-Mo-W alloy is about 800MPa at 1500˚C [21]. The strength of the Ir-based alloys is lower than that of this Nb-based alloy. A series of Ir-Ti, -Nb, -Hf, and -Zr alloys with a second-element concentration from 0 to 25at% were also compressed at room temperature, 1200˚C, and 1800˚C (Fig. 3). Here, the shadowed, solid, and open symbols show the compressive strength at room temperature, 1200˚C, and 1800˚C, respectively. The component phases in the alloys depend on the composition. For example, in the Ir-Nb alloy, the alloys have an fcc-single phase up to 11at% Nb. The L12 single phase exists above 24at% Nb. The fcc and L12 two-phase regions exist in alloys with concentrations between 11 and 24at% Nb. Although we investigated only one sample with an fcc single-phase because of sample limitations, the alloys containing a second element showed higher strength than pure Ir. Thus, we conclude that a solid-solution hardening effect was shown in Ir-based alloys. The strength of the fcc and L12 two-phase alloys was higher than that of the fcc or L12 single-phase ones at room temperature and 1200˚C in the case of Ir-Nb, -Hf, and -Zr alloys. We attribute the further strengthening achieved by adding a second element to the precipitation hardening of the L12 phase, as in Ni-based superalloys [7, 22]. In Ir-Ti alloys, the strength increment of the two-phase alloy was small. At 1800˚C, there was no clear precipitation-hardening difference among the alloys.
Fig. 3 Concentration dependence of the compressive strength of the Ir-(a) Ti, (b) Nb, (c) Hf, and (d) Zr alloys. The shadowed, solid, and open symbols show the compressive strength at room temperature, 1200˚C, and 1800˚C, respectively.
The contribution by solid-solution hardening can be estimated by the law of mixtures in the strength of composite materials as a linear combination of the strengths of the fcc matrix and the L12 precipitates. Assuming that the composition of the fcc and L12 phases in the alloy is at equilibrium, the strength of the fcc and L12 phases can be estimated as the strength of the fcc single phase and the L12 single phase with equivalent composition. We investigated the strength of an fcc single-phase alloy with a composition close to equilibrium value. For the L12 single phase, we investigated an alloy with only stoichiometric composition. We estimated the solid-solution hardening effect, as shown by the dotted line (for example, in the Ir-Zr alloy). The peak strength of the two-phase alloys was higher than the strength by solidsolution hardening. This difference between the experimentally determined strength and the strength by solid-solution hardening (arrows in Fig. 3d) is due to precipitation hardening. Precipitation hardening determined by the above method is shown as a function of the volume fraction of the precipitates in Fig. 4. Below 1200˚C, clear evidence for precipitation hardening was noted. The Ir-Zr alloy, with the maze structure generally showed higher precipitation hardening than other alloys at any testing temperature. The second highest effect as precipitation hardening was observed in the Ir- Fig. 4 Precipitation-hardening Nb alloys, which had an orderly arrayed function of volume fraction of the precipitates.
a
structure. The Ir-Hf alloy showed the third highest precipitation hardening. In the Ir-Ti alloy, precipitation hardening was not clear at any testing temperature. At 1800˚C, precipitation hardening was generally low and did not depend on the alloy although the Ir-Zr alloy showed slightly higher precipitation hardening than the other alloys. 3.3. DISLOCATION STRUCTURE The dislocation structure was investigated in the sample deformed until the plastic strain reached 2-3%. A real strain was investigated by measuring sample length before and after the test and is shown in the caption. Figures 5-7 show the dislocation structure of Ir-15Ti, -15Nb, 13Hf, and -15Zr alloys deformed at room temperature, 1200˚C, and 1800˚C. After deformation at room temperature, there was no clear microstructure difference between deformed and heattreated samples in the Ir-Ti and Ir-Nb alloys (Fig. 5a and b). In the dark-field image of the Ir-Ti alloy, the bright-contrast phase represents precipitates, in which dislocations were observed in them. This deformed microstructure is similar to the microstructure in the heat-treated sample (Fig. 1b). In the Ir-Nb alloy, a dislocation contrast was observed at the interface of the precipitates. We observed a similar contrast in a heat-treated sample under out-of-edge-on condition. Thus, we were unable to ascertain if these dislocations were formed during the deformation or the heat treatment. In the Ir-Hf and Ir-Zr alloy, we observed a drastic change. In the Ir-Hf alloy (Fig. 5c), the precipitate shows bright contrast. Precipitate shape was similar to that in the semi-coherent structure (Fig. 1g). However a lot of dislocations were observed in the deformed precipitates, and a dislocation tangle was observed at the interface between the matrix and the precipitates in the deformed sample. In the Ir-Zr alloy (Fig. 5d), the shape of the precipitate was similar to that of the semi-coherent structure and a dislocation tangle, similar to that observed at the interface in the Ir-Hf alloy, was also observed. This dislocation tangle was different from the misfit dislocations that aligned at the interface with regular spacing (Fig. 1h). Thus, we conclude that this dislocation tangle was formed during deformation. Shearing of precipitates on a plane leaving a stacking fault or anti-phase boundary was observed in a few places (arrows).
Fig. 5 Dislocation structure of the Ir-(a)15at%Ti, (b)15at%Nb, (c)13at%Hf, and (d)15at%Zr alloys. They were plastically deformed at room temperature (a) 0.6%, (b) 0.3%, (c) 0.7%, and (d) 0.1%, respectively. (a) Dark-field image. (b)-(d) Bright-field images. These images were −
−−
taken in the beam direction of (a) [001] (g=110 ), (b) [112] (g=220), (c) [101] (g=111 ), and (d) [001] (g=220).
After deformation at 1200˚C, no clear microstructure difference between deformed and heat-treated samples was observed in the Ir-Ti or the Ir-Nb alloys. As in the sample deformed at room temperature, dislocations were observed in the precipitates with bright contrast in the Ir-Ti alloy (Fig. 6a). Dislocation contrast was also observed at the interface in the Ir-Nb alloy (Fig. 6b). In the Ir-Hf alloy, dislocations became denser in the precipitates (bright contrast) and a dislocation tangle appeared at interface, compared with the structure of the sample deformed at room temperature (Fig. 6c). In the Ir-Zr alloy, dislocation tangles were observed such as the structure of the sample deformed at room temperature (Fig. 6d). Shearing by bowed dislocations in precipitates was observed, as shown by arrows. In addition, some dislocations, which had bowed into the precipitates from the interface, were also observed (A in Fig. 6d).
Fig. 6 Dislocation structure of the Ir-(a)15at%Ti, (b)15at%Nb, (c)13at%Hf, and (d)15at%Zr alloys. They were plastically deformed at 1200˚C (a) 0.5%, (b) 3.0%, (c) 0.5%, and (d) 3.2%, respectively. (a) Dark-field image. (b)-(d) Bright-field images. These images were taken in the beam direction of (a) [001] (g=110), (b) [112] (g=220), (c) [101] (g=020), and (d) [101] (g=020).
Figure 7 shows the deformation structure at 1800˚C. Here, we can see a drastic microstructure change in all of the alloys. Dark-field images were taken for the Ir-Ti and -Nb alloys. Precipitates are the bright-contrast phase. In the Ir-Ti alloy, the shape of the precipitates changed to spherical, and some precipitates became connected with each other. Dislocations were not observed in either the precipitates or the matrix (Fig. 7a). In the Ir-Nb alloy, some precipitates grew large or became connected with each other, reaching over 200 nm in size. Cuboidal precipitates changed to a rectangular shape (Fig. 7b). In the Ir-Hf alloy, the precipitates deformed and changed to irregular shapes. Many dislocations were observed in these precipitates (Fig. 7c). In the Ir-Zr alloy, we observed denser dislocation tangle at the interface, and more dislocations were observed in the precipitates, compared with the structure after deformation at room temperature (Fig. 7d). These results are summarized in Table 2.
Fig. 7 Dislocation structure of the Ir-(a)15at%Ti, (b)15at%Nb, (c)13at%Hf, and (d)15at%Zr alloys. They were plastically deformed at 1800˚C (a) 3.2%, (b) 0.3%, (c) 0.3%, and (d) 1.6%, respectively. (a)(b) Dark-field images. (c)(d) Bright-field images. These images were taken in −−
the beam direction of (a) [101] (g=010), (b) [101] (g=010), (c) [101] (g=111 ), and (d) [101] −−
(g=111).
Table 2 Dislocation structure after deformation Ir-Ti Ir-Nb Room Dislocations in Bypass temperature precipitates
1200˚C
Dislocations in precipitates
Bypass
1800˚C
Change of precipitate shape
Change of precipitate shape Coarsening
If-Hf Dislocation tangle
Ir-Zr Dislocation tangle
Shearing
Shearing on a plane Dislocation tangle
Dislocation tangle Shearing Change of precipitate shape Shearing
Shearing Shearing
4. DISCUSSION 4.1. DEFORMATION BEHAVIOR Our results have shown that precipitation hardening is more effective at large lattice misfit. Microstructure of the alloys depended on lattice misfit, as a result, deformation mode changed below 1200 ˚C. Dependence of lattice misfit, microstructure, and deformation mode on precipitation hardening is summarized in table 3. The Ir-Zr alloy with the maze structure formed by large lattice misfit showed highest precipitation hardening, although the Ir-Hf alloy with lattice misfit similar to the Ir-Zr alloy showed third highest value. The second strongest precipitation hardening was shown in the Ir-Nb alloy with an orderly arrayed structure formed by moderate lattice misfit. The Ir-Ti alloy with irregular shaped precipitates formed by small lattice misfit showed lowest precipitation hardening among the alloys tested. The deformation mode of the Ir-Zr and Ir-Hf alloys was found to be shearing precipitates. (Fig. 5c, d and 6d, d). Although we did not find clear evidence of the deformation mechanism of the Ir-Nb alloy by observation in TEM, we suggested bypass mechanisms in our previous study [15]. The deformation mechanism of the Ir-Ti alloy, with irregular precipitates was not observed clearly either. We observed dislocations in the precipitates in a heat-treated sample of the Ir-Ti alloy. This suggests that deformation of precipitates is easy if these dislocations move. Thus, we conclude that the deformation mechanism of the Ir-Ti alloy is also a shearing mechanism. These results showed that the deformation mechanism did not depend on lattice misfit unlike precipitation hardening. Table 3. Dependence of lattice misfit, microstructure, and deformation mode on precipitation hardening Precipitation Lattice misfit Microstructure Deformation mode hardening Ir-Zr Large Large (2%) Maze structure Shearing (Ir-Hf) Ir-Nb Moderate Moderate (0.3 ~ Orderly arrayed Bypass (Ir-Ta) 0.4%) structure Ir-Ti Small Small (-0.2 ~ Irregular structure Shearing (Ir-V) 0.1 %) Next, we discuss how the deformation mode and lattice misfit affected the magnitude of precipitation hardening. In the Ir-Zr alloy, a thin fcc matrix 10 nm in width surrounded by precipitates was observed. If slip plane of dislocation is supposed as {111} planes similar to general fcc alloys, a mixed structure of elongated trapezoids and triangles will appear on the {111} intersection in the maze structure [13]. Thus bypass of precipitates seems to be difficult for a dislocation in complex maze structure. On the other hand, bypass of precipitates seems to be easy in the Ir-Nb alloy because the width of the fcc matrix is between about 20 to 50 nm. Thus, a dislocation must shear the precipitates to deform in the Ir-Zr alloy. Dislocations moving in the fcc matrix often encounter precipitates because of the fineness of the microstructure in the Ir-Zr alloy. Furthermore, dislocations are trapped and form tangles at the coherent or semicoherent interface (Fig. 5d and 6d). When lattice misfit is large, coherency-strain field at the coherent interface will also be large and lot of misfit dislocations will form at the semi-coherent interface. These factors are effective for strong precipitation hardening of the Ir-Zr alloy. In the Ir-Nb alloy, a dislocation can move by bypassing precipitates without shearing. Thus the resistance for dislocation movement in the Ir-Nb alloys is lower than that of the Ir-Zr alloy. The precipitation hardening of the Ir-Ti alloy was lowest among the alloys. The large number of dislocations in the precipitates suggest that the precipitates and coherent interface are not effective for preventing movement of dislocation. We consider that when the lattice misfit is small, the prevention of dislocation movement at the coherent interfaces will also be small. When precipitates are not obstacles for a dislocation and the dislocation can move in both of the matrix and precipitates with small resistance, the dislocation does not need to bypass precipitates. For these two reasons, precipitation hardening of the Ir-Ti alloy is smaller than those of the IrZr and Ir-Nb alloys.
On the other hand, precipitation hardening of the Ir-Hf alloy was lower than that of the IrNb alloy although the lattice misfit of the Ir-Hf alloy is larger than that of the Ir-Nb alloy. This is because the macroscopic microstructure of the Ir-Hf alloy is heterogeneous. After the heat treatment, the dendrite structure shown in the as-cast sample did not disappear, and precipitates were formed only in the interdendritic region (Fig. 8). We believe that if homogeneous precipitation can be obtained in the Ir-Hf alloy, the precipitation hardening effect will be equivalent to that in the Ir-Zr alloy.
Fig. 8 Secondary electron images of Ir-15at%Hf heat-treated at 1400˚C for 168 hours. (b) is enlarged microstructure of (a).
Although precipitation hardening of the Ir-Ta alloy was not investigated, we consider that the precipitation-hardening effect and the deformation mechanism of the Ir-Ta alloy are the same as that of the Ir-Nb alloy because of their similar microstructures. We did not investigate the precipitation hardening of the Ir-V alloy because the Ir-V alloy is brittle and preparation of TEM samples is very difficult. The deformation mechanism of the Ir-V alloy may be different from that of the Ir-Ti alloy because of different precipitate morphology. This remains to be clarified. At 1800˚C, precipitation hardening was not effective in all the alloys. We observed that not only the matrix but also the precipitates deformed and changed their shape by deformation in IrTi, -Nb, and -Hf alloys. In the Ir-Zr alloy, shearing of the dislocation into precipitates often occurred. We concluded that both matrix and precipitate can deform, and the coherent interface is no longer effective to prevent the dislocation movement at 1800˚C. Thus, precipitation hardening was not effective. 4 . 2 . POTENTIAL OF IR-BASED ALLOYS AS HIGH-TEMPERATURE MATERIALS From the point of view of precipitation hardening, the Ir-Zr, -Nb, and -Ta alloys, which show high precipitation hardening, are promising as high-temperature materials. Coherent precipitates are suitable to obtain a high precipitation-hardening effect. In an Ir-based binary system, solution heat treatment is difficult because fcc solubility does not extend to the twophase region at high temperature. If fcc solubility can be expanded at high temperature by adding a third element and homogeneous precipitation can be obtained, the Ir-Hf alloy will also be a promising material. To use high-temperature materials, microstructure stability is an important factor. In Ir-Zr alloys, both coherent and semi-coherent structures were formed in the same sample. A discontinuous coarsened lamellar structure was also formed on grain boundaries and grew into the grains to release coherent strain energy [13]. This is due to the large lattice misfit and shows that the coherent structure is not stable in alloys with large lattice misfit compared with alloys with a small lattice misfit. If the sample with a large lattice misfit is heat-treated at high
temperature for a long time, the coherent structure will change to a semi-coherent structure, or a coarsened lamellar structure will grow in the whole area. Although the precipitation-hardening effect is strong for the alloy with a large lattice misfit, when the alloy is used at high temperature for long time, maintaining the coherent structure as long as possible is desirable. In fact, precipitation hardening of Ir-based alloys was not effective at 1800˚C. We must consider a stable coherent interface to prevent the dislocation movement at 1800˚C. The strength improvement of each phase, fcc or L12 , by solid solution strengthening is also important. Ma et al. showed that Nb-based alloys are strengthened by solid-solution hardening by W and Mo (refractory metal) as well as by precipitation hardening of silicide [21], obtaining high strength over 1500˚C. This suggests another way to bring about strengthening at high temperature. 5. CONCLUSIONS 1. 2.
3.
4.
An fcc and L12 two-phase coherent structure was observed by TEM in Ir-V, -Ti, -Nb, -Ta, -Hf, and -Zr alloys after heat treatment between 1200 and 1400 ˚C for up to 168 hours. The microstructure depends on the lattice misfit between the precipitates and the matrix. When the lattice misfit is from -0.2 ~ 0.1% (Ir-V and -Ti alloys) through 0.3-0.4% (Ir-Nb and -Ta alloys) to around 2% (Ir-Hf and -Zr alloys), the microstructure changes from irregular to orderly arrayed and maze structures. When the lattice misfit is large, a semicoherent structure also forms in addition to the coherent structure. Solid-solution hardening and precipitation hardening were observed in these alloys. Precipitation hardening depends on microstructure resulting from lattice misfit difference. The alloy with the maze structure showed highest precipitation hardening. Deformation mode of the alloy with the maze structure was shearing. Second highest precipitation hardening was shown in the alloy with an orderly arrayed structure, which deformed by the bypass mechanism. The lowest precipitation hardening was shown in the alloy with an irregular structure. The alloy was also deformed by shearing, however the coherent interface was not an effective obstacle to prevent dislocations because of small lattice misfit. At 1800˚C, precipitation-hardening effect was not clear because the coherent interface was not effective for prevention of dislocation movement.
ACKNOWLEDGMENTS We are grateful to Mr. S. Nishikawa of Furuya Metal Co., Ltd. for preparing the Ir-based alloy ingots.
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Y. Yamabe-Mitarai:
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