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Synthesis and thermal stability of nano-RuAl by mechanical alloying
Materials Science and Engineering A329– 331 (2002) 112– 117 www.elsevier.com/locate/msea
Synthesis and thermal stability of nano-RuAl by mechanical alloying K.W. Liu *, F. Mu¨cklich Department of Materials Science, Functional Materials, Saarland Uni6ersity, Geb. 22; Etage 7, P.O. Box 151150, D-66041 Saarbriu¨cken, Germany
Abstract Single phase RuAl has been synthesized by mechanical alloying (MA). The structural evolution during MA and thermal stability of as-milled powders at elevated temperatures have been analyzed by XRD, thermal analysis and isothermal annealing. The results indicate that there are two stages of reaction between Ru and Al in MA before single phase RuAl is obtained, and that the as-milled RuAl undergoes reordering, strain relaxation and grain growth at high temperatures. All structural evolutions show signs of stagnation upon high temperature annealing. The isothermal grain growth kinetics has been analyzed by considering the influence of impurities. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Nano-grained; Single phase RuAl; Mechanical alloying; Thermal stability; Grain growth kinetics
1. Introduction Due to its high melting point (about 2323 K), hightemperature strength, oxidation and corrosion resistance, and good room-temperature toughness, the intermetallic compound RuAl is regarded as a potential material for high temperature applications in aggressive environments [1,2]. The manufacturing methods employed to produce RuAl until now have generally included casting [3] and powder metallurgy (PM) route [4]. Both methods have failed to produce homogenous single phase RuAl. An alternative way of producing RuAl is the solid state reaction, namely mechanical alloying (MA) [5–7]. The solid state processing nature of MA makes it extremely suitable for producing high temperature intermetallics, especially the systems with constituent elements that have widely different melting points and that exist within a very narrow composition range. Several reports on nano-grained materials by MA showed promising thermal stability which may be related to the grain boundary segregation of solute atoms and/or dispersoid particles [8]. The nano-materials pro* Corresponding author. Tel.: +49-681-3023048; fax: + 49-6813024876. E-mail address: [email protected] (K.W. Liu).
duced by MA are prominently characterized by the existence of a large quantity of impurities introduced from either milling materials (Fe, Cr, or ceramic materials) or milling atmosphere (mainly O2 and N2) [9]. In view of the significant influence of solute/impurities on grain growth kinetics, Michels’s model [10] has been employed in this study to analyze the grain growth kinetics:
n
D 2(t)− D 2max 2kt = exp − 2 D max D 2(0)− D 2max
(1)
where D(0) is the initial mean grain size before annealing, D(t) the mean grain diameter at an annealing time t, Dmax maximum average grain size and k a constant. In the present work, single phase nano-grained RuAl is synthesized directly from powder mixture with a composition of Ru50Al50 (atomic percent) by mechanical alloying. Its thermal stability has been studied by thermal analysis and isothermal annealing.
2. Experimental Elemental powders of Ru with a purity of 99.94% and Al with a purity of 99.9% were used as the starting materials. The ball milling was performed in a laboratory Spex 8000 mill using hardened steel balls and a vial
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with a ball-to-powder weight ratio of 10:1. All handling of the powders during milling was performed in the glove box in an Argon atmosphere. The structural evolution of powders during milling was examined by X-ray diffraction (XRD) using CuKa radiation in a PHILIPS X’Pert X-ray diffractometer.
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Transmission electron microscopy (TEM) was performed on a JEOL 2000 microscope operated at 200 kV. Thermal analysis was carried out in a Netzsch DSC 404 at a heating rate of 20 K min − 1 in flowing argon gas. Isothermal annealing was performed by heating the powders at the fixed temperature for a certain time in the flow of Argon and H2 mixture. The level of iron impurity was examined by Energy-Dispersive X-ray Analysis (EDX) in EDX/REM JEOL 840A. The oxygen and nitrogen contamination were determined by hot extraction with an Evalograph VH9O.
3. Results
3.1. Structural e6olution during milling
Fig. 1. X-ray diffraction patterns for powders as a function of milling time.
Fig. 2. Impurity Fe content in powders as a function of milling time; the N2 and O2 content in powders milled for 35 h have also been attached on the graph.
The structural evolution of the powders during the milling time is followed by XRD (Fig. 1). The crystalline diffraction peaks of Ru and Al are evident in powders milled for 10 min. With the milling time extended up to 3 h 40 min, the diffraction peaks corresponding to Ru become broader and less intensity, while those corresponding to Al vanish from the XRD, indicating the dissolution of Al into the Ru matrix. The powders at this stage are referred to as Ru(Al) solid solution/mixture in the context. When the milling time is increased to 4 h, the XRD pattern shows that the powders consist of RuAl and a certain amount of Ru, which means that the original Ru(Al) solid solution/ mixture has been transformed into RuAl by an abrupt reaction during ball milling. The mixture of RuAl and Ru remains unchanged during a further milling time of up to 15 h. Finally, after milling for 25 h, no diffraction peak for elemental Ru and Al can be detected and the XRD result shows a single phase RuAl. From the above XRD results, grain sizes in powders are obtained by analyzing the broadening of diffraction peaks [11]. The calculated results show the grain size decreases rapidly from the starting size to less than 30 nm after 4 h of milling. The grain size of RuAl resulting from the reaction at the milling time of 4 h is about 12 nm. With a further increase of the milling time, the grain size of RuAl decreases quite sluggishly from 12 to about 5 nm after 25 h of milling and remains, however, nearly constant from a milling time of 25–50 h. The nature of the phases formed and the corresponding grain sizes determined by XRD in powders during milling have been further verified by TEM, for detail see [12]. The level of iron impurity is shown in Fig. 2 as a function of the milling time. The oxygen and nitrogen contamination for powders milled for 35 h is also attached on the graph. The Fe content remains almost unchanged from the initial value in elemental powders with less than 4 h of milling. With a prolonged milling
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Fig. 3. DSC scans for powders milled for different times.
Structural evolution for powders milled for different time at elevated temperatures was further identified by isothermal annealing treatment of powders milled for different times. Characteristic temperatures (573, 873 and additional 1073 K) just before and after the exothermic peak have been chosen as annealing temperatures. For powders milled for 4 h, the phase constitution of RuAl and Ru remains unchanged after annealing at 873 and 1073 K for 0.5 h. For powders milled for 15 h, a trace of Ru still remains in powders after annealing at 873 K, but has been transformed into RuAl by annealing at 1073 K. The final product after annealing at 1073 K for 0.5 h is a single phase RuAl. The structural evolution for powders milled for 35 and 50 h are quite similar. In light of the re-appearance of reflection corresponding to superlattice (100), the disordered single phase RuAl after annealing at 573 K becomes ordered after annealing at 873 K for 0.5 h. The grain size is in the range of 5–7 nm for powders after annealing at 573 K, 10–18 nm at 873 K and 20–40 nm at 1073 K. The lattice strain decreases from 1.2 to 2% for powders annealed at 573 K to around 0.5% for those annealed at 873 K and finally to 0.2% at 1073 K. The grain sizes and phases in powders milled for different times and followed by annealing treatments have been verified by TEM. The grain size of about 20 nm and a single phase RuAl obtained by TEM in powders milled for 35 h and followed by annealing at 1073 K for 0.5 h (Fig. 4) agree with those obtained by XRD.
3.3. Isothermal grain growth kinetics
Fig. 4. Dark field image and corresponding diffraction pattern for powders milled for 35 h with subsequent annealing at 800 °C for 0.5 h.
time from 4 to 50 h, the Fe content increases approximately linearly with milling time from the initial value to about 35 at.%.
3.2. Structural e6olution of as-milled powders at ele6ated temperatures The DSC results for powders milled for different durations are shown in Fig. 3. From the results, several exothermic peaks can be seen at a temperature lower than 1006 K for powders milled for 2 h and 3 h 40 min. The exothermic releases are believed to be generated by the formation of RuAl from the Ru(Al) solid solution/ mixture. On the one hand, there is almost no sign of heat release for powders milled for 4, 15 and 25 h, on the other, an exothermic peak is observed to start from about 662 K and end at about 785.2 K for powders milled for 35 and 50 h.
As-milled single phase RuAl (MA 35 h) with a grain size of about 5 nm has been used as a starting material for isothermal grain growth studies. Isothermal annealing has been performed at 873, 973, 1073, 1173 and 1273 K for 0.5, 1, 2, 3 and 5 h. The overall results indicate that no phase transformation of the as-milled RuAl has been detected except variation of diffraction sharpening at different annealing temperatures and times. The evaluated grain sizes at different annealing temperatures as a function of annealing time are shown in Fig. 5. Grain growth stagnates after certain time of annealing at all temperatures, for instance, after 1 h at 873 and 973 K. At each temperature, the data of grain sizes were fitted into Eq. (1) to determine best fit values for the parameters k and Dmax. The dashed lines in Fig. 5 reflect the fitted curves. The measured grain sizes at all temperatures are well fitted by Eq. (1) as can be seen in Fig. 5. The activation energy E can be subsequently obtained from the slope of a plot of lnk against 1/T [11] (Fig. 6). Apparently, the overall points do not fit to a straight line, although three points corresponding 873, 973 and 1073 K do. It is better to fit the data points
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with a curve instead. I have tried to get the corresponding activation energy from the slope of neighboring points for each temperature range. The apparent grain growth activation energies evaluated in this way are listed in Table 1. One can see that the grain growth activation energies E increase dramatically with higher annealing temperature ranges. At a fixed annealing temperature, internal strains
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drop down abruptly from 1.7% (as-milled) to values measured after 0.5 h of annealing, then decay slowly with increasing annealing time, and finally reach a level that seems not to change upon prolonged treatment time. The long-range (LR) order parameter S/S0 was determined from the intensities of the superlattice in relation to the fundamental XRD lines of the CsCl structure. At a certain annealing temperature, for example 1073 K, LR ordering reaches a level at the shortest annealing time (0.5 h) and then remains quite stable with further extended annealing time (up to 5 h).
4. Discussion
4.1. Formation mechanism for RuAl during milling
Fig. 5. Grain size as a function of annealing at different temperatures, points are measured data and dotted lines are fitted with Eq. (1).
Fig. 6. Arrhenius plot of the parameter k (in Eq. (1)) against the reciprocal of annealing temperatures. Table 1 Grain growth activation energies at different temperature ranges Temperature range (K)
Activation energy E (kJ mol−1)
873–1073 1073–1173 1173–1273
39.0 72.2 213.5
The high ductility of Al makes it easy to de deformed and stick on the surfaces of the container and milling balls, and, therefore, the amount of Al participating in the alloying process decreases to a lower level after the first 4 h of milling. This results in a mixture of RuAl and Ru after the reaction producing RuAl. Further milling deforms Al on the surfaces of container and milling balls extensively and makes them fracture and fall into the container gradually to join the alloying process. This alloying process leads to powders with phase constituents of RuAl and Ru(Al) solid solution/ mixture. The process ends when RuAl is produced from the Ru(Al) solid solution/mixture by another abrupt reaction during milling (MA 25 h). Further milling is exerted on single phase RuAl. There are, therefore, two stages of reaction and alloying processes before the single phase RuAl is formed during mechanical alloying of the Ru and Al powder mixture under the present milling conditions. The first stage (I) is the formation of the Ru(Al) solid solution/mixture followed by an abrupt reaction consuming the available Ru and Al elements in powders to form RuAl. No Al remains after the reaction. The second stage (II) is the formation of the Ru(Al) solid solution/mixture again by consumption of the remaining Ru and Al coated on the surfaces of the container and on the milling balls. The Ru(Al) solid solution/mixture is then transformed into RuAl by prolonged milling. The duration of the second stage depends on the rate at which Al is released from the coated surfaces. The final product of milling is a single phase RuAl. Further structural evolution after the second stage is mainly the disordering of the formed RuAl by ball impact. The single phase RuAl is the only intermetallic compound produced either by exposing the powders milled for shorter times (2 h) to a high temperature (\973 K) or by directly milling to a prolonged time (25 h). The
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Fig. 7. Variation of lattice parameter as a function of grain size.
process is quite significant because single phase RuAl with the same composition is difficult to obtain either by casting [3] or SHS processes [4]. The chemical uniformity of the constitutional atoms produced by mechanical alloying is believed to provide a basis for forming intermetallic compound corresponding to equilibrium phase diagram. Such an abrupt reaction is thought to be stimulated by a mechanism similar to self propagating high temperature synthesis, which is ignited by mechanical energy from impinging balls instead of by thermal energy [13]. It has been suggested that a large exothermic heat release during reaction between the components may be a prerequisite preceding the reaction. The formation enthalpy of RuAl, − 124 9 3.3 kJ mol − 1 [14], is comparable with −118.5 kJ mol − 1 for NiAl [15], −137 kJ mol − 1 for NbSi2 [16] and − 129 kJ mol − 1 for MoSi2 [17], respectively. Once the reaction is ignited by milling in one region, a large amount of heat release from the reaction will increase the temperature of neighboring regions where the reaction becomes self-sustaining.
4.2. Thermal stability of as-milled RuAl at ele6ated temperatures The overall results of annealing treatments for powders milled for different time indicate that the RuAl forming reaction during mechanical alloying is irreversible at high temperatures. More interesting is that all three kinds of structural evolutions (reordering, strain relaxation and grain growth) show some stagnation phenomena during isothermal annealing. The good representation of the grain growth data for all temperatures by Eq. (1) indicates the existence of a grain-size dependant pinning effect on the moving grain boundaries. The total amount of Fe has been detected to be as high as about 15 at.% in as-milled RuAl by EDX, while
the concentration of other possible impurities, such as Cr and O2 is detected to be negligible low, lower than 1 at.% (Fig. 2). Considering the fact that no diffraction peak corresponding to a second phase because RuAl has been detected from the XRD results and that TEM has failed to reveal any second phase of any kind in the matrix, the impurities might exist in the form of RuAl(Fe) solid solution and/or segregate at grain boundaries. Similar to Co and Ni which can be used as substitutional alloying elements for Ru in RuAl alloys because of the existence of B2 CoAl and NiAl phases [18], Fe could also act as a substitutional element for Ru in RuAl due to the B2 structure of FeAl. The isothermal section of the Fe–Ru –Al system at 820 K [19,20] unfolds a solubility up to 15 at.% Fe in RuAl. Considering that atomic size for Ru (1.89 A, ) is larger than that for Fe (1.72 A, ), the lattice constant of RuAl(Fe) solid solution would be expected to be smaller than that of pure RuAl if part of the Ru atoms were replaced by Fe. The lattice parameter of our RuAl samples has been calculated using the Cohen method [21] and is plotted as a function of grain size in Fig. 7. The lattice constants of all annealed RuAl samples are apparently smaller than the value of 2.9916 A, of pure RuAl [1]. The increment of the lattice parameter of our RuAl(Fe) solid solution with increasing grain size indicates the diffusion of Fe atoms out of the RuAl matrix. It is, therefore, believed that Fe atoms exist in the as-milled RuAl powders in the form of RuAl(Fe) solid solution. During grain growth the segregation of impurities in grain boundaries is expected to occur by considering the following reasons. First, the lattice constant increases almost linearly with increasing grain size at the beginning until the grain size reaches a value of about 40 nm. Then the lattice constant remains nearly constant with further grain growth (Fig. 7). Second, no second phase was found from XRD in samples after annealing at 1273 K for 5 h in XRD. The increase of the grain growth activation energy with increasing temperature is thought to be closely related to the accumulation of impurities in grain boundaries during grain growth. The increase of the grain growth activation energy with solute additives can also be observed by comparing the activation energy in single phase Al3Nb with that of the same phase alloying with Ti. The grain growth activation energy of Al3Nb has been obtained to be 16292 kJ mol − 1 [22] whilst that of Al3Nb – 5.5Ti is 2019 2 kJ mol − 1 [23]. Regarding the present case, similarly, the concentration of impurities in grain boundaries increases with the coarsening of the stagnated grain sizes at different temperatures and, therefore, higher activation energy for grain growth could be possible with rising temperatures. The evolution of the degree of ordering and the level of lattice strain during grain growth would be an
K.W. Liu, F. Mu¨ cklich / Materials Science and Engineering A329–331 (2002) 112–117
advantage to the grain growth stagnation by interaction with grain growth kinetics. The complex interaction between ordering and recrystallization reactions has been reported by Burkley [24] and Rajkovic and Burkley [25] to occur in the recrystallization of FeCo containing small amount of V or Cr. In their reports homogenous or continuous ordering increases rapidly between 873 and 1000 K with an ordering temperature of about 1000 K, so that the deformed alloy becomes ordered before the onset of recrystallization. The driving force for recrystallization is thus reduced substantially and the mobility of the interfaces involved in the recrystallization process may also be severely reduced by the ordering of the deformed matrix [24]. Similarly, the strain relaxation process during grain growth takes place by moving the atoms in distorted positions to perfect lattice positions, the process accompanied by a reduction in further diffusion which means a certain part of driving force for grain growth is taken away as well. Therefore, the occurrence of reordering and strain relaxation processes accompanying grain growth would contribute to grain growth stagnation by taking part of grain growth driving force away.
5. Conclusions Single phase RuAl can be obtained directly from an abrupt reaction during mechanical alloying of powder mixture of Ru and Al. There are two stages of alloying and reaction between Ru and Al during MA before single phase RuAl is obtained under the present milling conditions. All Al in powders has been consumed to form RuAl in the first stage and results in a mixture of RuAl and Ru. Further alloying of Al with Ru depends on the release of coated Al from the surfaces of the container/balls and ends up with another reaction to form RuAl. The structural evolutions of as-milled single phase RuAl upon high temperature exposures exhibit reordering, strain relaxation and grain growth. All three kinds of structural evolutions show signs of stagnation. The total amount of 15 at.% Fe has been found to exist in the form of solid solution in RuAl as substitutional element and as segregation in grain boundaries. The lattice diffusion and segregation of impurity atoms in grain boundaries during grain growth have been verified by the variation of lattice parameters. Reordering and strain relaxation processes accompanying grain growth would result in structure relaxation and thus take away a certain part of the driving force for grain growth. The grain size less than 40 nm annealed at 1173
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K for as long as 5 h indicates the strong thermal stability of the nano-RuAl.
Acknowledgements The authors would like to thank Dr C.E. Krill, A. Michels and Dr K. Smidoda for stimulating discussions, and J. Schmauch for DSC and M. Schule for TEM measurements. Scholarship to K.W. Liu from DFG Graduiertenkolleg (Saarbru¨ icken) is acknowledged. The research work of Professor Mu¨ cklich was supported by the Alfried Krupp Prize for Young University Teachers of the Krupp Foundation.
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R.L. Fleischer, Metall. Trans. 24A (1993) 227. I.M. Wolff, JOM 49 (1997) 34. I.M. Wolff, G. Sauthoff, Metall. Trans. 27A (1996) 2642. I.M. Wolff, Metall. Trans. 27A (1996) 3688. J.S. Benjamin, Sci. Am. 234 (5) (1976) 40. C.C. Koch, Nanostruct. Mater. 2 (1993) 109. K.W. Liu, J.S. Zhang, J.G. Wang, G.L. Chen, J. Mater. Res. 13 (1998) 1198. C.C. Koch, Nanostruct. Mater. 2 (1993) 109. B.S. Murty, S. Ranganatham, Int. Mater. Rev. 43 (1998) 101. A. Michels, C.E. Krill, H. Ehrhardt, R. Birringer, D.T. Wu, Acta Mater. 47 (1999) 2143. H.P. Klug, L. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, Second ed, Wiley, NY, 1974, p. 618. K.W. Liu, F. Mu¨ cklich, Intermetallics, 9 (2001) 81. S.N. Patankar, S.Q. Xiao, J.J. Lewandowski, A.H. Heuer, J. Mater. Res. 8 (1993) 1311. W.G. Jung, O.J. Kleppa, Metall. Trans. 23B (1992) 53. C.J. Smithells (Ed.), Metals Reference Book, Fifth ed., Butterworths, London, 1976, p. 193. A.K. Niessen, F.R. de Boer, R. Boom, P.F. de Chatel, W.C.M. Mattens, A.R. Miedema, CALPHAD 7 (1983) 51. B.K. Yen, T. Aizawa, J. Kihara, Mater. Sci. Eng. A220 (1996) 8. I.M. Wolff, G. Sauthoff, L.A. Cornish, H.D. Steyn, R. Coetzee, in: M.V. Nathal, R. Darolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics, TMS, Warrandale, PA, 1997, p. 815. V.F. Tsurikov, E.M. Sokolovskaya, T.P. Loboda, Izv, Akad. Nauk. SSSR, Metally, 6 (1980a), 226. G. Petzow, G. Effenberg (Eds.), Ternary Alloys. A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams, vol. 4, VCH, Weinheim, 1991, p. 251. B.D. Cullity, Elements of X-Ray Diffraction, Addison-Wesley, Reading, MA, 1978. J. Isonishi, K. Okazaki, J. Mater. Sci. 28 (1993) 3829. S. Kawanishi, K. Isonishi, K. Okazaki, Mater. Trans. JIM 34 (1993) 49. R.A. Buckley, Metal Sci., 1979, 67. M. Rajkovic, R.A. Buckley, Metal Sci., 1981, 21.
Synthesis and thermal stability of nano-RuAl by mechanical alloying K.W. Liu *, F. Mu¨cklich Department of Materials Science, Functional Materials, Saarland Uni6ersity, Geb. 22; Etage 7, P.O. Box 151150, D-66041 Saarbriu¨cken, Germany
Abstract Single phase RuAl has been synthesized by mechanical alloying (MA). The structural evolution during MA and thermal stability of as-milled powders at elevated temperatures have been analyzed by XRD, thermal analysis and isothermal annealing. The results indicate that there are two stages of reaction between Ru and Al in MA before single phase RuAl is obtained, and that the as-milled RuAl undergoes reordering, strain relaxation and grain growth at high temperatures. All structural evolutions show signs of stagnation upon high temperature annealing. The isothermal grain growth kinetics has been analyzed by considering the influence of impurities. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Nano-grained; Single phase RuAl; Mechanical alloying; Thermal stability; Grain growth kinetics
1. Introduction Due to its high melting point (about 2323 K), hightemperature strength, oxidation and corrosion resistance, and good room-temperature toughness, the intermetallic compound RuAl is regarded as a potential material for high temperature applications in aggressive environments [1,2]. The manufacturing methods employed to produce RuAl until now have generally included casting [3] and powder metallurgy (PM) route [4]. Both methods have failed to produce homogenous single phase RuAl. An alternative way of producing RuAl is the solid state reaction, namely mechanical alloying (MA) [5–7]. The solid state processing nature of MA makes it extremely suitable for producing high temperature intermetallics, especially the systems with constituent elements that have widely different melting points and that exist within a very narrow composition range. Several reports on nano-grained materials by MA showed promising thermal stability which may be related to the grain boundary segregation of solute atoms and/or dispersoid particles [8]. The nano-materials pro* Corresponding author. Tel.: +49-681-3023048; fax: + 49-6813024876. E-mail address: [email protected] (K.W. Liu).
duced by MA are prominently characterized by the existence of a large quantity of impurities introduced from either milling materials (Fe, Cr, or ceramic materials) or milling atmosphere (mainly O2 and N2) [9]. In view of the significant influence of solute/impurities on grain growth kinetics, Michels’s model [10] has been employed in this study to analyze the grain growth kinetics:
n
D 2(t)− D 2max 2kt = exp − 2 D max D 2(0)− D 2max
(1)
where D(0) is the initial mean grain size before annealing, D(t) the mean grain diameter at an annealing time t, Dmax maximum average grain size and k a constant. In the present work, single phase nano-grained RuAl is synthesized directly from powder mixture with a composition of Ru50Al50 (atomic percent) by mechanical alloying. Its thermal stability has been studied by thermal analysis and isothermal annealing.
2. Experimental Elemental powders of Ru with a purity of 99.94% and Al with a purity of 99.9% were used as the starting materials. The ball milling was performed in a laboratory Spex 8000 mill using hardened steel balls and a vial
0921-5093/02/$ - see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 5 4 1 - 6
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with a ball-to-powder weight ratio of 10:1. All handling of the powders during milling was performed in the glove box in an Argon atmosphere. The structural evolution of powders during milling was examined by X-ray diffraction (XRD) using CuKa radiation in a PHILIPS X’Pert X-ray diffractometer.
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Transmission electron microscopy (TEM) was performed on a JEOL 2000 microscope operated at 200 kV. Thermal analysis was carried out in a Netzsch DSC 404 at a heating rate of 20 K min − 1 in flowing argon gas. Isothermal annealing was performed by heating the powders at the fixed temperature for a certain time in the flow of Argon and H2 mixture. The level of iron impurity was examined by Energy-Dispersive X-ray Analysis (EDX) in EDX/REM JEOL 840A. The oxygen and nitrogen contamination were determined by hot extraction with an Evalograph VH9O.
3. Results
3.1. Structural e6olution during milling
Fig. 1. X-ray diffraction patterns for powders as a function of milling time.
Fig. 2. Impurity Fe content in powders as a function of milling time; the N2 and O2 content in powders milled for 35 h have also been attached on the graph.
The structural evolution of the powders during the milling time is followed by XRD (Fig. 1). The crystalline diffraction peaks of Ru and Al are evident in powders milled for 10 min. With the milling time extended up to 3 h 40 min, the diffraction peaks corresponding to Ru become broader and less intensity, while those corresponding to Al vanish from the XRD, indicating the dissolution of Al into the Ru matrix. The powders at this stage are referred to as Ru(Al) solid solution/mixture in the context. When the milling time is increased to 4 h, the XRD pattern shows that the powders consist of RuAl and a certain amount of Ru, which means that the original Ru(Al) solid solution/ mixture has been transformed into RuAl by an abrupt reaction during ball milling. The mixture of RuAl and Ru remains unchanged during a further milling time of up to 15 h. Finally, after milling for 25 h, no diffraction peak for elemental Ru and Al can be detected and the XRD result shows a single phase RuAl. From the above XRD results, grain sizes in powders are obtained by analyzing the broadening of diffraction peaks [11]. The calculated results show the grain size decreases rapidly from the starting size to less than 30 nm after 4 h of milling. The grain size of RuAl resulting from the reaction at the milling time of 4 h is about 12 nm. With a further increase of the milling time, the grain size of RuAl decreases quite sluggishly from 12 to about 5 nm after 25 h of milling and remains, however, nearly constant from a milling time of 25–50 h. The nature of the phases formed and the corresponding grain sizes determined by XRD in powders during milling have been further verified by TEM, for detail see [12]. The level of iron impurity is shown in Fig. 2 as a function of the milling time. The oxygen and nitrogen contamination for powders milled for 35 h is also attached on the graph. The Fe content remains almost unchanged from the initial value in elemental powders with less than 4 h of milling. With a prolonged milling
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Fig. 3. DSC scans for powders milled for different times.
Structural evolution for powders milled for different time at elevated temperatures was further identified by isothermal annealing treatment of powders milled for different times. Characteristic temperatures (573, 873 and additional 1073 K) just before and after the exothermic peak have been chosen as annealing temperatures. For powders milled for 4 h, the phase constitution of RuAl and Ru remains unchanged after annealing at 873 and 1073 K for 0.5 h. For powders milled for 15 h, a trace of Ru still remains in powders after annealing at 873 K, but has been transformed into RuAl by annealing at 1073 K. The final product after annealing at 1073 K for 0.5 h is a single phase RuAl. The structural evolution for powders milled for 35 and 50 h are quite similar. In light of the re-appearance of reflection corresponding to superlattice (100), the disordered single phase RuAl after annealing at 573 K becomes ordered after annealing at 873 K for 0.5 h. The grain size is in the range of 5–7 nm for powders after annealing at 573 K, 10–18 nm at 873 K and 20–40 nm at 1073 K. The lattice strain decreases from 1.2 to 2% for powders annealed at 573 K to around 0.5% for those annealed at 873 K and finally to 0.2% at 1073 K. The grain sizes and phases in powders milled for different times and followed by annealing treatments have been verified by TEM. The grain size of about 20 nm and a single phase RuAl obtained by TEM in powders milled for 35 h and followed by annealing at 1073 K for 0.5 h (Fig. 4) agree with those obtained by XRD.
3.3. Isothermal grain growth kinetics
Fig. 4. Dark field image and corresponding diffraction pattern for powders milled for 35 h with subsequent annealing at 800 °C for 0.5 h.
time from 4 to 50 h, the Fe content increases approximately linearly with milling time from the initial value to about 35 at.%.
3.2. Structural e6olution of as-milled powders at ele6ated temperatures The DSC results for powders milled for different durations are shown in Fig. 3. From the results, several exothermic peaks can be seen at a temperature lower than 1006 K for powders milled for 2 h and 3 h 40 min. The exothermic releases are believed to be generated by the formation of RuAl from the Ru(Al) solid solution/ mixture. On the one hand, there is almost no sign of heat release for powders milled for 4, 15 and 25 h, on the other, an exothermic peak is observed to start from about 662 K and end at about 785.2 K for powders milled for 35 and 50 h.
As-milled single phase RuAl (MA 35 h) with a grain size of about 5 nm has been used as a starting material for isothermal grain growth studies. Isothermal annealing has been performed at 873, 973, 1073, 1173 and 1273 K for 0.5, 1, 2, 3 and 5 h. The overall results indicate that no phase transformation of the as-milled RuAl has been detected except variation of diffraction sharpening at different annealing temperatures and times. The evaluated grain sizes at different annealing temperatures as a function of annealing time are shown in Fig. 5. Grain growth stagnates after certain time of annealing at all temperatures, for instance, after 1 h at 873 and 973 K. At each temperature, the data of grain sizes were fitted into Eq. (1) to determine best fit values for the parameters k and Dmax. The dashed lines in Fig. 5 reflect the fitted curves. The measured grain sizes at all temperatures are well fitted by Eq. (1) as can be seen in Fig. 5. The activation energy E can be subsequently obtained from the slope of a plot of lnk against 1/T [11] (Fig. 6). Apparently, the overall points do not fit to a straight line, although three points corresponding 873, 973 and 1073 K do. It is better to fit the data points
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with a curve instead. I have tried to get the corresponding activation energy from the slope of neighboring points for each temperature range. The apparent grain growth activation energies evaluated in this way are listed in Table 1. One can see that the grain growth activation energies E increase dramatically with higher annealing temperature ranges. At a fixed annealing temperature, internal strains
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drop down abruptly from 1.7% (as-milled) to values measured after 0.5 h of annealing, then decay slowly with increasing annealing time, and finally reach a level that seems not to change upon prolonged treatment time. The long-range (LR) order parameter S/S0 was determined from the intensities of the superlattice in relation to the fundamental XRD lines of the CsCl structure. At a certain annealing temperature, for example 1073 K, LR ordering reaches a level at the shortest annealing time (0.5 h) and then remains quite stable with further extended annealing time (up to 5 h).
4. Discussion
4.1. Formation mechanism for RuAl during milling
Fig. 5. Grain size as a function of annealing at different temperatures, points are measured data and dotted lines are fitted with Eq. (1).
Fig. 6. Arrhenius plot of the parameter k (in Eq. (1)) against the reciprocal of annealing temperatures. Table 1 Grain growth activation energies at different temperature ranges Temperature range (K)
Activation energy E (kJ mol−1)
873–1073 1073–1173 1173–1273
39.0 72.2 213.5
The high ductility of Al makes it easy to de deformed and stick on the surfaces of the container and milling balls, and, therefore, the amount of Al participating in the alloying process decreases to a lower level after the first 4 h of milling. This results in a mixture of RuAl and Ru after the reaction producing RuAl. Further milling deforms Al on the surfaces of container and milling balls extensively and makes them fracture and fall into the container gradually to join the alloying process. This alloying process leads to powders with phase constituents of RuAl and Ru(Al) solid solution/ mixture. The process ends when RuAl is produced from the Ru(Al) solid solution/mixture by another abrupt reaction during milling (MA 25 h). Further milling is exerted on single phase RuAl. There are, therefore, two stages of reaction and alloying processes before the single phase RuAl is formed during mechanical alloying of the Ru and Al powder mixture under the present milling conditions. The first stage (I) is the formation of the Ru(Al) solid solution/mixture followed by an abrupt reaction consuming the available Ru and Al elements in powders to form RuAl. No Al remains after the reaction. The second stage (II) is the formation of the Ru(Al) solid solution/mixture again by consumption of the remaining Ru and Al coated on the surfaces of the container and on the milling balls. The Ru(Al) solid solution/mixture is then transformed into RuAl by prolonged milling. The duration of the second stage depends on the rate at which Al is released from the coated surfaces. The final product of milling is a single phase RuAl. Further structural evolution after the second stage is mainly the disordering of the formed RuAl by ball impact. The single phase RuAl is the only intermetallic compound produced either by exposing the powders milled for shorter times (2 h) to a high temperature (\973 K) or by directly milling to a prolonged time (25 h). The
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Fig. 7. Variation of lattice parameter as a function of grain size.
process is quite significant because single phase RuAl with the same composition is difficult to obtain either by casting [3] or SHS processes [4]. The chemical uniformity of the constitutional atoms produced by mechanical alloying is believed to provide a basis for forming intermetallic compound corresponding to equilibrium phase diagram. Such an abrupt reaction is thought to be stimulated by a mechanism similar to self propagating high temperature synthesis, which is ignited by mechanical energy from impinging balls instead of by thermal energy [13]. It has been suggested that a large exothermic heat release during reaction between the components may be a prerequisite preceding the reaction. The formation enthalpy of RuAl, − 124 9 3.3 kJ mol − 1 [14], is comparable with −118.5 kJ mol − 1 for NiAl [15], −137 kJ mol − 1 for NbSi2 [16] and − 129 kJ mol − 1 for MoSi2 [17], respectively. Once the reaction is ignited by milling in one region, a large amount of heat release from the reaction will increase the temperature of neighboring regions where the reaction becomes self-sustaining.
4.2. Thermal stability of as-milled RuAl at ele6ated temperatures The overall results of annealing treatments for powders milled for different time indicate that the RuAl forming reaction during mechanical alloying is irreversible at high temperatures. More interesting is that all three kinds of structural evolutions (reordering, strain relaxation and grain growth) show some stagnation phenomena during isothermal annealing. The good representation of the grain growth data for all temperatures by Eq. (1) indicates the existence of a grain-size dependant pinning effect on the moving grain boundaries. The total amount of Fe has been detected to be as high as about 15 at.% in as-milled RuAl by EDX, while
the concentration of other possible impurities, such as Cr and O2 is detected to be negligible low, lower than 1 at.% (Fig. 2). Considering the fact that no diffraction peak corresponding to a second phase because RuAl has been detected from the XRD results and that TEM has failed to reveal any second phase of any kind in the matrix, the impurities might exist in the form of RuAl(Fe) solid solution and/or segregate at grain boundaries. Similar to Co and Ni which can be used as substitutional alloying elements for Ru in RuAl alloys because of the existence of B2 CoAl and NiAl phases [18], Fe could also act as a substitutional element for Ru in RuAl due to the B2 structure of FeAl. The isothermal section of the Fe–Ru –Al system at 820 K [19,20] unfolds a solubility up to 15 at.% Fe in RuAl. Considering that atomic size for Ru (1.89 A, ) is larger than that for Fe (1.72 A, ), the lattice constant of RuAl(Fe) solid solution would be expected to be smaller than that of pure RuAl if part of the Ru atoms were replaced by Fe. The lattice parameter of our RuAl samples has been calculated using the Cohen method [21] and is plotted as a function of grain size in Fig. 7. The lattice constants of all annealed RuAl samples are apparently smaller than the value of 2.9916 A, of pure RuAl [1]. The increment of the lattice parameter of our RuAl(Fe) solid solution with increasing grain size indicates the diffusion of Fe atoms out of the RuAl matrix. It is, therefore, believed that Fe atoms exist in the as-milled RuAl powders in the form of RuAl(Fe) solid solution. During grain growth the segregation of impurities in grain boundaries is expected to occur by considering the following reasons. First, the lattice constant increases almost linearly with increasing grain size at the beginning until the grain size reaches a value of about 40 nm. Then the lattice constant remains nearly constant with further grain growth (Fig. 7). Second, no second phase was found from XRD in samples after annealing at 1273 K for 5 h in XRD. The increase of the grain growth activation energy with increasing temperature is thought to be closely related to the accumulation of impurities in grain boundaries during grain growth. The increase of the grain growth activation energy with solute additives can also be observed by comparing the activation energy in single phase Al3Nb with that of the same phase alloying with Ti. The grain growth activation energy of Al3Nb has been obtained to be 16292 kJ mol − 1 [22] whilst that of Al3Nb – 5.5Ti is 2019 2 kJ mol − 1 [23]. Regarding the present case, similarly, the concentration of impurities in grain boundaries increases with the coarsening of the stagnated grain sizes at different temperatures and, therefore, higher activation energy for grain growth could be possible with rising temperatures. The evolution of the degree of ordering and the level of lattice strain during grain growth would be an
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advantage to the grain growth stagnation by interaction with grain growth kinetics. The complex interaction between ordering and recrystallization reactions has been reported by Burkley [24] and Rajkovic and Burkley [25] to occur in the recrystallization of FeCo containing small amount of V or Cr. In their reports homogenous or continuous ordering increases rapidly between 873 and 1000 K with an ordering temperature of about 1000 K, so that the deformed alloy becomes ordered before the onset of recrystallization. The driving force for recrystallization is thus reduced substantially and the mobility of the interfaces involved in the recrystallization process may also be severely reduced by the ordering of the deformed matrix [24]. Similarly, the strain relaxation process during grain growth takes place by moving the atoms in distorted positions to perfect lattice positions, the process accompanied by a reduction in further diffusion which means a certain part of driving force for grain growth is taken away as well. Therefore, the occurrence of reordering and strain relaxation processes accompanying grain growth would contribute to grain growth stagnation by taking part of grain growth driving force away.
5. Conclusions Single phase RuAl can be obtained directly from an abrupt reaction during mechanical alloying of powder mixture of Ru and Al. There are two stages of alloying and reaction between Ru and Al during MA before single phase RuAl is obtained under the present milling conditions. All Al in powders has been consumed to form RuAl in the first stage and results in a mixture of RuAl and Ru. Further alloying of Al with Ru depends on the release of coated Al from the surfaces of the container/balls and ends up with another reaction to form RuAl. The structural evolutions of as-milled single phase RuAl upon high temperature exposures exhibit reordering, strain relaxation and grain growth. All three kinds of structural evolutions show signs of stagnation. The total amount of 15 at.% Fe has been found to exist in the form of solid solution in RuAl as substitutional element and as segregation in grain boundaries. The lattice diffusion and segregation of impurity atoms in grain boundaries during grain growth have been verified by the variation of lattice parameters. Reordering and strain relaxation processes accompanying grain growth would result in structure relaxation and thus take away a certain part of the driving force for grain growth. The grain size less than 40 nm annealed at 1173
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K for as long as 5 h indicates the strong thermal stability of the nano-RuAl.
Acknowledgements The authors would like to thank Dr C.E. Krill, A. Michels and Dr K. Smidoda for stimulating discussions, and J. Schmauch for DSC and M. Schule for TEM measurements. Scholarship to K.W. Liu from DFG Graduiertenkolleg (Saarbru¨ icken) is acknowledged. The research work of Professor Mu¨ cklich was supported by the Alfried Krupp Prize for Young University Teachers of the Krupp Foundation.
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