Effects of boron on the microstructure and thermal stability of directionally solidified NiAl–Mo eutectic

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Acta Materialia 58 (2010) 421–428 www.elsevier.com/locate/actamat

Effects of boron on the microstructure and thermal stability of directionally solidified NiAl–Mo eutectic A. Gali a,b, H. Bei a, E.P. George a,b,* a

Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6115, USA b Department of Materials Science and Engineering, University of Tennessee, Knoxville, TN 37996-2200, USA Received 15 May 2009; received in revised form 31 August 2009; accepted 6 September 2009 Available online 7 October 2009

Abstract Microalloying with 0.01 at.% B decreases the range of growth speeds over which a well-aligned fibrous eutectic microstructure can be obtained in directionally solidified NiAl–Mo. Compared to the undoped alloy, the size/spacing of the Mo fibers is larger, and the fiber density smaller, in the B-doped alloy. Annealing at 1400 °C coarsens the fibers by a mechanism involving fault migration and annihilation driven by diffusion along the fiber–matrix interface. The coarsening kinetics, given by the decrease in Mo fiber density with time, is exponential, and microalloying with B decreases the coarsening rate. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Interfaces; Segregation; Directional solidification; Coarsening; Composites

1. Introduction The NiAl–Mo eutectic has the composition Ni–45.5Al– 9Mo (all compositions in this paper are in at.% unless otherwise noted) and solidifies as a fibrous composite consisting of long Mo-alloy fibers embedded in an NiAl matrix. It has received considerable attention [1–10] because of its attractive properties for high-temperature structural applications, including a high melting point (1600 °C) and better toughness [2], creep strength [9] and tensile strength [10] compared to the monolithic intermetallic compound NiAl. However, its microstructural stability at elevated temperatures, which can influence mechanical properties, has not been extensively investigated. The purpose of this paper is to report the kinetics of microstructural degradation in this alloy and the effects of boron added as a microalloying element.

* Corresponding author. Address: Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6115, USA. E-mail address: [email protected] (E.P. George).

Microstructural degradation in multi-phase alloys can be manifested as a size and/or a shape change of the constituent phases as discussed recently [11] for another eutectic alloy, Cr–Cr3Si, which has a lamellar rather than fibrous microstructure. It is driven by the minimization of total interfacial free energy per unit volume, GI, which, in the absence of stresses, can be written as: X GI ¼ Ai ci ð1Þ i

where Ai is the total interfacial area per unit volume of interface i, and ci is the interfacial energy per unit area of that interface. Minimization of GI can occur as a result of decreases in Ai and/or ci. In fibrous composites such as NiAl–Mo, the total interfacial area can be decreased by increasing the average fiber size (which is equivalent to increasing the average interfiber spacing for constant fiber volume fraction). One way of achieving this is by decreasing the solidification rate. The other term in Eq. (1), interfacial energy per unit area (ci), can be decreased by segregation of alloying elements to the fiber–matrix interface. In this study, we report the effects of the microalloying element boron, which was

1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.09.020

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chosen because of its limited solubility in both Mo [12] and NiAl [13]. In general, the tendency for an element to segregate to the grain boundaries varies inversely as its solubility in the bulk [14,15]. Relatively small amounts of B (0.01%) were added to minimize the likelihood of formation of a third phase in the microstructure. 2. Experimental procedures Alloys with compositions Ni–45.5Al–9Mo (undoped NiAl–Mo) and Ni–45.445Al–9Mo–0.01B (B-doped NiAl– Mo) were arc-melted 25 times in argon atmosphere to ensure thorough mixing. Thereafter, they were drop cast into a cylindrical copper mould 10 mm in diameter and 120 mm in length. The starting materials (>99.5% pure by weight) were carefully weighed before melting. Total weight losses after drop casting were found to be small, <0.1%. Therefore, all the final compositions are assumed to be the same as the aim compositions. Directional solidification was performed as discussed previously [10] in a xenon-arc-lamp floating-zone furnace in flowing argon gas. The drop-cast rods were used as feed material, and pieces cut from directionally solidified material were used as seeds. To obtain a homogeneous molten zone, the seed and feed rods were rotated in opposite directions at 60 rpm. A stable molten zone was maintained by adjusting the power of the lamp as needed. To study the effect of B on the as-grown microstructure, the B-doped alloy was directionally solidified at speeds to 60 mm h1. Based on these results, a growth speed of 40 mm h1 was chosen to investigate subsequent microstructural stability in the undoped and B-doped alloys. Specimens measuring 2.5 mm in thickness were cut from the directionally solidified alloys by electric discharge machining, ground with SiC paper through 600 grit, and annealed in vacuum (2  105 torr) at 1400 °C for various times. The annealed specimens were mounted in epoxy, ground again through 600 grit, and polished on a “vibromet” polishing machine through 0.3 lm Al2O3. The polished specimens were etched using a solution of 80% HCl and 20% HNO3 and examined in an FEI XL 30 FEG scanning electron microscope (SEM). Representative scanning electron micrographs were analyzed using image analysis software to measure the fiber density, fiber size and area fraction of fibers. Orientation relationships between the Mo fibers and NiAl matrix were determined by electron back scattered diffraction patterns acquired in the FEI SEM. 3. Results and discussion 3.1. Microstructure of directionally solidified B-doped NiAl– Mo The undoped NiAl–Mo alloy investigated in the present study exhibits a fibrous microstructure similar to that which has been reported in the literature [6,10,16]. The

effect of B addition on this microstructure was studied by directional solidification at different growth speeds. Fig. 1 shows the microstructures of B-doped NiAl–Mo grown at 20, 30, 50 and 60 mm h1. Adding 0.01% B to NiAl– Mo decreases the range of growth speeds over which a fibrous microstructure can be obtained. A well-aligned fibrous microstructure was obtained only for speeds up to 50 mm h1. At the higher growth rate of 60 mm h1, the B-doped alloy exhibits a cellular microstructure (Fig. 1d), suggesting that a transition to cellular solidification occurs between 50 and 60 mm h1 for the rotation speed of 60 rpm investigated here. By contrast, in the undoped alloy a fibrous microstructure can be obtained for growth speeds as high as 160 mm h1 [17]. These observations are consistent with what we know about regular eutectic microstructures, namely, that: (i) they tend to be obtained when the solid–liquid interface is planar and stable [18,19] and (ii) alloying elements added to a eutectic system (e.g. B to NiAl–Mo) tend to destabilize the planar solid–liquid interface resulting in cellular microstructures [20]. Figs. 2 and 3 show the fiber spacing (k, the average distance on a transverse section between the centers of adjacent Mo fibers) and fiber size (a, the average edge length of approximately square cross-sections), respectively, as a function of the inverse square root of the growth rate (R). In both figures, best fits to the data were obtained for straight lines, similar to that obtained previously for the undoped NiAl–Mo alloy [10]. This is in accordance with the prediction made by Jackson and Hunt [21]. According to their theory, under maximum growth rate conditions,   DnC 0 M ra rb ð2Þ þ k2 R ¼ 1 2ð1 þ nÞ2 mb mb where ra and rb are the solid–liquid interfacial energies for the constituent phases a and b [20], M is a tabulated function of volume fraction [22], D is the solution diffusion coefficient, C0 is the alloy composition and ma, mb are the slopes of the liquidus lines corresponding to the a and b phases, respectively. In fibrous eutectics where the fibers are arranged in a hexagonal pattern, which is approximately the case for NiAl–Mo [10], the parameter n ¼ a=ðk  aÞ can be shown to be a constant. Therefore, Eq. (2) can be rewritten as: k2 R ¼ constant

ð3Þ

Additionally, if the fiber volume fraction is constant, a similar relationship can be derived [10] between fiber size (a) and growth rate (R), namely: a2 R ¼ constant

ð4Þ

Therefore, plots of k vs. R1/2 and a vs. R1/2 are expected to be straight lines, as shown in Figs. 2 and 3. For a given growth rate, the Mo fibers in the B-doped alloy are spaced farther apart (Fig. 2) and are larger in size (Fig. 3) than those in the undoped alloy. However, the

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Fig. 1. SEM micrographs showing the microstructure of B-doped NiAl–Mo: (a) 20 mm h1; (b) 30 mm h1; (c) 50 mm h1 and (d) 60 mm h1. Wellaligned fibrous structure was obtained to 50 mm h1. Cellular structure was formed at 60 mm h1. The inset shows the polygonal morphology of Mo fibers on the transverse section.

growth directions of the fibers and matrix ([1 0 0]NiAl k [1 0 0]Mo) are unaffected by the addition of boron, as are the cross-sectional morphologies of the fibers. In general, the fibers in the B-doped and undoped alloys are polygonal in shape, but often they have approximately square (inset, Fig. 1a) or rectangular cross-sections, both of which can exhibit cropped corners sometimes.

Fig. 2. Effect of growth rate on the average inter-fiber spacing of Mo fibers in B-doped NiAl–Mo and NiAl–Mo. Data for the NiAl–Mo alloy were taken from Ref. [10].

3.2. Thermal stability of directionally solidified NiAl–Mo and B-doped NiAl–Mo To determine the effect of B on microstructural stability, directionally solidified alloys grown at 40 mm h1 and 60 rpm were annealed at 1400 °C for times to 400 h. The resulting microstructures are shown in Figs. 4 and 5. The volume fraction of the Mo fibers in the directionally solid-

Fig. 3. Effect of growth rate on the fiber size in B-doped NiAl–Mo and NiAl–Mo. Data for the NiAl–Mo alloy were taken from Ref. [10].

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Fig. 4. Microstructure of directionally solidified NiAl–Mo: (a) as-grown and (b) after annealing for 400 h at 1400 °C.

ified alloys, as determined from the measured area fractions, is 14% in both the undoped and B-doped alloys. The average size of the Mo fibers increased, and the number of Mo fibers decreased, with increasing annealing time in both alloys. Fig. 6 shows a plot of the fiber density (number of fibers intersecting a transverse section per unit area) as a function of annealing time. Best fit was obtained for an exponential curve as predicted by Weatherly and Nakagawa [23] for coarsening by fault migration and annihilation in fibrous eutectic alloys. Note that, since a steady state was not observed for the fiber size distribution during annealing, the Lifshitz and Slyozov [24] and Wagner [25] equations cannot be used to determine the coarsening rate of our Mo fibers. Schematic diagrams showing growth faults and the fault migration and annihilation mechanism are provided in Fig. 7. Terminations are formed in fibrous eutectics during directional solidification to maintain a constant inter-fiber spacing under steady-state growth conditions [26]. The chemical potential of terminations is higher than that of branches because of their greater total curvature [27]. As a result, material transport can occur from the terminations to branches. Material transport can also occur from terminations to neighboring fibers that have smaller curvatures. In both cases, the fault length (dis-

Fig. 5. Microstructure of directionally solidified B-doped NiAl–Mo: (a) as-grown and (b) after annealing for 400 h at 1400 °C.

Fig. 6. Fiber density as a function of annealing times in NiAl–Mo and Bdoped NiAl–Mo eutectic alloys.

tance between a branch and its termination) decreases with time until the termination disappears (Fig. 7b and c). This fault annihilation mechanism causes the number of rods that intersect a transverse section to decrease with time.

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Fig. 7. Schematic illustration of coarsening mechanism in fiber eutectics by fault migration and annihilation.

At any given time, the fiber density N (number of fibers intersecting a transverse section per unit area) can be written as [23]: N ¼ N T þ nL

ð5Þ

where NT is the density of continuous fibers (i.e. fibers that run through the directionally solidified rod continuously), n is the number of faults per unit volume and L is the average fault length. If the faults are randomly distributed along the growth direction, the fiber density remains the same along the growth direction for a given growth speed. In that case, the rate at which N decreases is given by: @N @L N  N T @n ¼n þ n @t @t @t

ð6Þ

The first term on the right is due to the flux of migrating faults, while the second term is due to annihilation of faults. As a first approximation [27], the average migration ¼ kN , where k is a conrate of faults can be written as @L @t stant [27]. Eq. (6) then becomes @N N  N T @n ¼ kNn þ : n @t @t

ð7Þ

The general solution for Eq. (7) is an exponential function, but an analytical solution exists only for particular cases [23]. Case I: all faults have the same length. For time t < s, where s is the time taken for the faults to ¼ 0 and n = n0 (number of faults per unit volannihilate, @n @t ume in the initial microstructure). Then Eq. (7) can be written as: @N ¼ kn0 N @t which yields the following solution N ¼ N 0 expðkn0 tÞ

ð8Þ

The apparent decrease in fiber density observed on the transverse section, for t < s, is due to the flux of the migrating faults. Case II: variable fault lengths in the initial microstructure.

The rate of decrease of the number of faults per unit volume is given by [27]: @n n ¼ @t s

ð9Þ

where s ¼ L1 @L (average time taken for the faults to @t annihilate). Eq. (7) can then be written as: @N N  NT ¼ kNn þ s @t

ð10Þ

If the number of faults per unit volume is much larger than the number of continuous fibers in the microstructure, i.e. N  NT, Eq. (10) becomes   dN 1 ¼ kn þ dt ð11Þ N s If s is a constant, the solution for Eq. (9), n = n0 exp(t/s), can be inserted into Eq. (11) to obtain: n t o ð12Þ N ¼ N 0 exp   kn0 sð1  expðt=sÞÞ s [27] and If s is a variable then s ¼ 1 nk N 2 ¼ ð1 þ ktn0 Þ N0

ð13Þ

We will now examine whether either of these two special cases applies to our results. Figs. 8–10 show the longitudinal microstructures of the directionally solidified undoped and B-doped alloys. It can be seen that faults of varying length are present in the initial microstructure (Fig. 8). Arrows 1 in Fig. 9 show Mo fibers that have been cut into two, probably because the longitudinal section is not through the fiber midsection and is not perfectly parallel to the fiber axis. Arrows 2 show those faults that have become thinner and tapered along the length after annealing. This tapered shape is similar to the simulated profiles attained by a solid of revolution by surface diffusion [28] and consistent with the notion that mass transport occurs by interfacial diffusion (since the solubility of Mo in the NiAl matrix is practically zero [10], it is difficult to get significant Mo transport through the

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Fig. 8. SEM micrographs showing microstructures of longitudinal sections of: (a) NiAl–Mo and (b) B-doped NiAl–Mo directionally solidified at 40 mm h–1 and 60 rpm. Faults of varying lengths are present in each of these alloys.

matrix). The mechanism of coarsening is therefore fault migration and annihilation. After annealing at 1400 °C for 400 h the fiber density decreases by less than an order of magnitude and the coarsening curve starts to flatten (Fig. 6). Therefore, the condition N  NT is not satisfied. This inequality is more likely to be satisfied in fibrous composites that are solidified at very high growth rates and exhibit the “Rayleigh instability” [29], such as Al–Al3Ni [30–32], where the fibers break up into a row of small spheres due to longitudinal perturbations and NT decreases. In the NiAl–Mo eutectics, we saw no evidence of the Rayleigh instability (which has to be caught in the early stages as longitudinal perturbations on the fibers before fully isolated spheres are formed that can be washed away during the etching). This could be due to the anisotropic energy of the Mo–NiAl interfaces (which is evident from the polygonal shapes of the Mo fiber cross-sections). When the interfacial energy is anisotropic (i.e. the Wulff plot has cusps), perturbations are unstable and tend to flatten with time. Together, the above observations lead us to conclude that neither of the particular cases discussed above applies to our NiAl–Mo system. Therefore the coarsening kinetics (Fig. 6) can only be represented as an empirically determined exponential curve, which, as mentioned earlier, is the general form of the solution to Eq. (7). The decrease in fiber density in the NiAl–Mo alloy after 400 h (60% of its initial value) is similar to that obtained

Fig. 9. SEM micrographs showing longitudinal sections of directionally solidified NiAl–Mo alloy after annealing at 1400 °C for: (a) 324 h and (b) 400 h. Arrows 1 show fibers that were cut into two and arrows 2 show tapered faults.

by Walter and Cline [16] for similar growth conditions (50%). Unlike us, however, they concluded that no significant coarsening occurred in this system. A possible reason for these different interpretations is that the data points obtained here are spread more uniformly along the time scale whereas those obtained by Walter and Cline were mostly concentrated at times below 100 h, with only two data points between 100 and 400 h. Also, care was taken by us to measure the fiber density in areas where the cooling rates are identical: since the fiber density depends on the cooling rate, regions which cool faster (e.g. outer surfaces of directionally solidified rods) will have a higher density than regions which cool more slowly (e.g. centers). Additional confirmation of the exponential coarsening kinetics can be obtained by assuming, following Ardell [33], that the volume fraction of fibers remains constant during annealing: N ðtÞAðtÞ ¼ N 0 A0 ¼ constant

ð14Þ

where N(t) and A(t) are the values of the fiber density and average cross-sectional area of the fibers after annealing for time t, and N0 and A0 are the corresponding values in the

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inversely, should increase exponentially. Since area can be measured independently of fiber density, the above equation can be used to confirm the shape of the coarsening curve. Fig. 11 shows how the average cross-sectional area of fibers varies as a function of annealing time. For both alloys, best fit again was obtained for an exponential curve consistent with the curve fit shown in Fig. 6. Adding 0.01% B decreases the coarsening rate of NiAl– Mo (Fig. 6) but does not change the fiber volume fraction. There are several possible reasons that could explain the observed B-induced decrease in the coarsening rate: (i) a decrease in the thermodynamic driving force for coarsening, GI, in the initial microstructure, because the initial fiber density, and hence the total surface area, is smaller in the B-doped alloy (as shown in Fig. 6); (ii) the fault density in the B-doped alloy is smaller than that in the undoped alloy (because n0k = constant [26] and, for a given growth rate, kNiAl–Mo < kNiAl–Mo+B); (iii) the average fault length, and hence s, is greater in the B-doped alloy; and (iv) boron changes the energy and/or structure of the NiAl–Mo interfaces, thereby making interface diffusion and migration harder. The present work has shown the plausibility of (i) and (ii). Additional work is needed to investigate possibilities (iii) and (iv). It is certainly reasonable to expect a decrease in the interface energy if B segregates to the NiAl–Mo interfaces. Therefore, future work should be aimed at detecting this segregation, for example, by chemical analysis in a high-resolution transmission electron microscope. Fig. 10. SEM micrographs showing longitudinal sections of directionally solidified B-doped NiAl–Mo alloy after annealing at 1400 °C for: (a) 196 h and (b) 400 h. Arrows 2 show tapered faults.

Fig. 11. Average area of fibers as a function of annealing times in NiAl– Mo and B-doped NiAl–Mo eutectic alloy.

as-solidified condition. Thus, if the fiber density decreases exponentially with time, the average area, which behaves

4. Conclusions The microstructure and thermal stability of NiAl–Mo eutectics, with and without 0.01% B added as a microalloying element, were investigated. Boron decreases the range of growth speeds for which a well-aligned fibrous microstructure can be obtained. At a rotation rate of 60 rpm, cellular microstructures were obtained at a growth rate of 60 mm h1, whereas well-aligned fibrous microstructures were obtained at growth rates of 50 mm h1 or less. The size/spacing of the Mo fibers is larger, and the fiber density smaller, in the B-doped alloy compared to the undoped NiAl–Mo alloy. The cross-sectional morphology of the Mo fibers (polygonal) is not affected by the addition of boron. Coarsening of the undoped and B-doped NiAl– Mo alloys was studied at 1400 °C for times up to 400 h and was found to occur by fault migration and annihilation. Fiber densities decrease exponentially with time and the coarsening rate of the B-doped alloy is lower than that of the undoped NiAl–Mo alloy. Possible reasons for the slower kinetics were discussed. Molybdenum fibers are stable with respect to longitudinal perturbations and do not exhibit the Rayleigh instability. Tapering of the faults suggests that the diffusion that drives the observed coarsening occurs along the fiber–matrix interfaces rather than through the matrix, which is consistent with the very low solubility of Mo in the NiAl matrix.

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Acknowledgements Research sponsored by the Division of Materials Sciences and Engineering, U.S. Department of Energy. A portion of this research was conducted at the SHaRE User Facility, which is sponsored by the Office of Basic Energy Sciences, Division of Scientific User Facilities, U.S. Department of Energy. References [1] Pickard SM, Zhang H, Ghosh AK. Acta Mater 1997;45:4333. [2] Misra A, Wu ZL, Kush MT, Gibala R. Philos Mag A 1998;78:533. [3] Heredia FE, He MY, Lucas GE, Evans AG, Deve HE, Konitzer D. Acta Metall Mater 1993;41:505. [4] Misra A, Gibala R. Intermetallics 2000;8:1025. [5] Raynolds JE, Roddick ER, Smith JR, Srolovitz DJ. Acta Mater 1999;47:3281. [6] Ferrandini P, Batista WW, Caram R. J Alloys Compd 2004;381:91. [7] Sundar RS, Kitazono K, Sato E, Kuribayashi K. Acta Mater 2001;49:1717. [8] Fox MR, Ghosh AK. Mater Sci Eng A 1999;A259:261. [9] Johnson DR, Oliver BF, Noebe RD, Whittenberger JD. Intermetallics 1995;3:493. [10] Bei H, George EP. Acta Mater 2005;53:69.

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