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TiAl–Nb melt interaction with pyrolytic boron nitride crucibles
Materials Chemistry and Physics 119 (2010) 347–350
Contents lists available at ScienceDirect
Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Materials science communication
TiAl–Nb melt interaction with pyrolytic boron nitride crucibles A.V. Kartavykh a,∗ , V.V. Tcherdyntsev b , J. Zollinger c,1 a
Institute of Chemical Problems for Microelectronics (IChPM), B. Tolmachevsky per. 5, 119017 Moscow, Russia Moscow State Institute of Steel and Alloys (Technological University), Leninsky pr. 4, 119049 Moscow, Russia c ACCESS e.V., Intzestrasse 5, 52072 Aachen, Germany b
a r t i c l e
i n f o
Article history: Received 31 March 2009 Received in revised form 10 September 2009 Accepted 18 September 2009 Keywords: Intermetallic compounds Solidification Ceramics Surface properties
a b s t r a c t This paper studies the chemical interaction on contact, and the related evolution of composition and microstructure of cast TiAl-based alloy, when using refractory oxygen-free pyrolytic boron nitride (pBN) crucibles/moulds instead of traditional oxide ceramic ones. We extend the research begun with AlN crucibles and reported in [4]. Three melting tests were performed in resistive SiC electro-furnace with Ti–46Al–8Nb (at%) alloy in 6N-purity pBN boat crucibles with fixed melt superheating times of 5, 12 and 25 min at 1670 ◦ C and subsequent quenching with high-purity Ar gas flow. As-cast samples were examined by optical microscopy/metallography, X-ray diffractometry, SEM-BSE, EDX and EBSD microanalyses, and compared to the initial material. Unlike the AlN it was shown that dissolved BN reacts with the melt, forming Ti2 AlN and (Ti,Nb)B micro-precipitates at the expense of selective destruction of the ␣2 -Ti3 Al major intermetallic phase in the solidified alloy: Ti3 Al + (Nb) + BN → Ti2 AlN + (Ti,Nb)B. The key features and temporal dynamics of this interaction are studied. © 2009 Elsevier B.V. All rights reserved.
1. Introduction TiAl-based alloys are good candidate materials for aero engine and gas-burning power-generating turbine blades [1]. One of the best processing routes for ␥-TiAl-based alloys is near-net shape investment casting in ceramic moulds in a high-purity inert gas atmosphere. The choice of mould material is a critical stage of the technological process. No refractories which are absolutely chemically inert against TiAl melts are known today [2,3]. The high solidification temperatures (1550–1600 ◦ C depending on alloying) and the melt reactivity both require more thermodynamically stable and chemically inert materials for mould manufacturing, instead of traditionally applied zirconia (ZrO2 ) and more advanced yttria (Y2 O3 -based) ceramics. One of the main reasons why the properties and structure of alloy being solidified are hard to reproduce is contamination with interstitial oxygen, which is difficult to control. The problem, mechanisms and consequences of uncontrolled oxygen contamination of TiAl intermetallics from crucibles/moulds at the melt processing have been observed elsewhere [4]. Thus, testing of alternative non-oxide (oxygenless) refractory materials may help in finding crucible materials which are better suited for processing the TiAl-based alloys. The results of such an experimental study, performed with the aluminium nitride
∗ Corresponding author. Tel.: +7 499 7889063; fax: +7 495 9538869. E-mail address: [email protected] (A.V. Kartavykh). 1 Present address: LETAM, Université Paul Verlaine de Metz, Ile du Saulcy, F-57012 Metz, Cedex, France. 0254-0584/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2009.09.021
crucibles, have previously been reported [4]. It was shown that AlN apparently holds promise as a mould material usable in metallurgical processing of TiAl intermetallics for the reduction of interstitial oxygen content. The nearest analogue to AlN in the series of refractory ceramics-like nitrides is boron nitride, having similar chemical and even higher thermal stability. This paper presents our investigations of the chemical compatibility of BN as the casting mould material for Ti–46Al–8Nb alloy. 2. Experimental procedures 2.1. Materials The pBN boat crucibles used in this work are serial commercial BoralloyTM products supplied from GE Advanced Ceramics, USA, having a volume of 2 cm3 , a mass of 1.5 g and wall thickness of 1.5 mm. They are manufactured by gas phase multilayer deposition, providing extra high-purity material of 6N grade [5]. The pyrolytic BN is stable up to 2500 ◦ C in argon; does not contain chemically bonded oxygen; it has zero open porosity, rather high thermal conductivity up to 60 W m−1 K−1 and linear thermal expansion coefficient of (2.2–4.4) × 10−6 K−1 promoting excellent resistance to thermal shock [5,6]. The pyrolytic BN is widely used in crucibles for processing of various intermetallics, alloys and semiconductor materials (see, for example [7,8]). Of particular importance is the fact that pBN maintains its resistance to corrosive attack by aluminum even at extreme temperature. Besides, boron (one of the constituents of the crucible material) is a routine grain refiner in as-cast TiAl-based alloys and is specifically used sometimes as a minor addition in the melt, at levels below 1 at.% [9]. Dissolved nitrogen (another constituent of crucible) does not have any detrimental effect on the microstructure of TiAl–Nb alloys, at least at contents less than 1800 wt ppm [4]. All the abovementioned prerequisites were taken into account in the choice of pBN as the possible alternative mould material for testing. The intermetallic was supplied through the IMPRESS project. Ti–46Al–8Nb (nominal composition, in at.%) alloy was prepared with induction melting and subsequent triple homogenizing vacuum arc remelting. Samples for tests have the cylindrical shape Ø 15 mm × 15 mm. A batch of 10 samples was delivered with the
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Table 1 Compatibility test conditions and XRD results for the initial Ti–46Al–8Nb material and as-cast samples. Sample number
Superheating time at 1670 ◦ C in pBN crucible (min)
Initial material No. 1 No. 2 No. 3
0 5 12 25
Phases revealed in the sample’s bulk and their volumetric fractions (%)
TiAl
Ti3 Al
Ti2 AlN
85.1 ± 0.2 84.8 ± 0.2 72.0 ± 0.5 78.9 ± 0.2
14.9 ± 0.2 15.2 ± 0.2 22.0 ± 0.5 0
0 0 6.0 ± 0.5 21.1 ± 0.2
Structural type of phase
tP2/1
hP8/3
hP8/4
Lattice periods (Å)
A = 2.841 ± 0.002 C = 4.084 ± 0.002
A = 5.780 ± 0.004 C = 4.654 ± 0.004
A = 3.003 ± 0.004 C = 13.628 ± 0.004
analysis certificate giving the real composition of Ti–46.35Al–7.87Nb (at.%). Our vacuum fusion analysis has established O, N, C and S interstitial impurities content of 700, 40, 230 and 25 wt ppm, respectively, in the alloy batch. 2.2. Experimental apparatus and conditions The experiments were performed in a special low-inertial electro-furnace, specified in detail in paper [4]. The facility consists of a vertical hermetic quartz flow-through reactor, connected to a gas-purifying line through the bottom flange. The reactor blow-through proceeds upward. Inside the chamber, a flat SiC/graphitecomposite resistive heater is installed, on which a boat crucible containing the test alloy sample may be directly placed. The sample temperature is measured by optical pyrometer with ±20 ◦ C accuracy and governed by heater power control. The small volume and mass of the test alloy charge (5 g) and crucible (1.5 g) both ensure the heat treatment of the sample occurs in near bulk-isothermal conditions. Each experiment was performed with the standard quartz tube-case being exchanged for a new one, to avoid any pollution accumulation in the reactor and to reduce the chemical background effect on material contamination level. During the experiment, technological argon was used. Ar passed through triple refining: drying in silica gel column following O2 and CO2 absorption in zeolite filters. The applied gas purification technique is standard for the semiconductor industry; its quality was controlled by measuring the dew point temperature, which did not exceed −70 ◦ C. Before each experiment, the charged furnace was evacuated down to ∼10−2 Torr pressure and then intensively blasted with Ar stream for 30 min in the cold state. Test cycles contained the following stages: heating of the crucible/sample pair up to 1670 ◦ C (i.e., 100 ◦ C in excess of the equilibrium liquidus of the test alloy [10]) at a rate of 120 ◦ C min−1 ; isothermal superheating of 5, 12 or 25 min duration, defining the fixed time of melt/crucible interaction; rapid cooling (quenching) at a rate of 15 ◦ C s−1 to preserve the interaction results. An Ar flow rate of 0.03 L s−1 at 120 kPa pressure was applied during the melting and superheating stages. Quenching was performed by upward blowing over the heater with a strong argon stream (five times increased Ar flow rate) while simultaneously switching off the power. The correspondence between the sample numbers and their superheating parameters is given in Table 1.
of the sample superheated for 5 min. The sample surface is mat black with apparently some BN scales on it—a first indication of an interaction between alloy and crucible. Fig. 1(b) shows that this mat black crust also exists on the free surface. The transverse vertical cut made reveals that the material has a metallic appearance, and also indicates good wetting between the melt and crucible from the meniscus shape of the sample free surface.
2.3. Analytical techniques Analytical techniques were applied to metallographic slices representing the vertical transverse cross-sections of samples (mini-ingots) (see Fig. 1). The 1-m diamond polishing of slices was done using Buehler equipment and consumables; no chemical etching was applied. Optical microscopy was performed using a POLYVAR Reichert–Jung microscope, with main the focus on identifying reaction layers in the vicinity of the former interface of the melt/crucible, and precipitates in the bulk of TiAl-based alloy. All reaction products revealed were subjected afterwards to electron-probe micro-analysis. The local phase composition was measured using scanning electron microscopy (SEM) and energy-dispersive X-ray spectrometry (EDX) techniques. For backscattered electron (BSE) imaging and for EDX a Gemini 1550 microscope was used with the INCA-Energy software. The crystal structure of precipitates was checked by electron backscatter diffraction (EBSD) using the crossbeam microscope 1540 XB and the software HKL Flamenco. For X-ray phase analysis (XRD), a DRON-3 diffractometer with CuK␣ radiation source was used, equipped with a specialized holder for study of bulk textured specimens with the angle of sample positioning at 30◦ . For averaging of the output diffraction, the rotation of each fixed TiAl–Nb slice was applied with respect to its normal. Spectra interpretation was carried out using software [11] allowing the analysis of relative phase content by the Rietveld method, with ±0.2% accuracy.
3. Results and discussion Fig. 1 gives a typical overview of as-processed sample. In Fig. 1(a) part of the pBN crucible is shown next to the crucible-adjacent side
Fig. 1. Photographs of (a) BN crucible and sample re-solidified after 5 min of the melt superheating (No. 1 in Table 1); (b) as-cut metallographic slice of the same sample.
A.V. Kartavykh et al. / Materials Chemistry and Physics 119 (2010) 347–350
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Fig. 2. Optical micrographs (on the left) and SEM-BSE microstructure images (on the right) at the interface between the reaction layer and Ti–46Al–8Nb alloy in samples processed at 1670 ◦ C in pBN crucible with (a and b) 5 min, (c and d) 12 min, and (e and f) 25 min superheating time. No definite boundary was established between the reaction layer and the alloy bulk in the latter sample whose whole volume is uniformly filled with needle-like precipitates.
3.1. Interaction layers and microstructure of cast alloy Microstructural evolution of the alloy as dependent on superheating time is depicted in Fig. 2. Pronounced interaction layers were found by optical microscopy in the samples superheated at relatively short durations, and clearly seen in Fig. 2(a and c). The transient zone between the reaction layer and the main bulk is shown in the samples held for 5 and 12 min in Fig. 2(b and d), using SEM-BSE mapping. The reaction layer can be observed at the bottom of these images and is composed of a black matrix with some white needles. On the upper part of these images, one observes the solidification microstructure of the Ti–46Al–8Nb alloy appearing in light grey contrast in the BSE mode. This microstructure is composed of the two intermetallic ordered phases ␣2 -Ti3 Al (D019 ) and ␥-TiAl (L10 ) that forms a lamellar texture. It can clearly be seen that black particles have formed ahead of the reaction layer front, whose volume fraction increases with superheating time. Moreover, in Fig. 2(d), some randomly distributed white needles appear in the whole microstructure. The microstructure of the sample after 25 min superheating is shown in Fig. 2(e and f) at equal magnifi-
cations. It can be seen from optical microscopy/metallography in Fig. 2(e) that the whole sample bulk is filled with short-needlelike precipitates. Its microstructure is composed of black and white particles in Fig. 2(f), and the residual alloy itself (appearing in grey BSE contrast in this image) has lost the lamellar structure and is composed now of a single phase. 3.2. Phase composition and evolution XRD analyses were performed on the bulk part of the as-cast samples as well as on the initial material. The results are shown in Table 1. The following events were observed with increasing superheating time: (i) after 12 min, the ␥/␣2 ratio decreased with respect to the initial material, i.e., the volume fraction of ␣2 increased. The as-cast alloy also contains up to 6% volumetric fraction of Ti2 AlN phase; (ii) after 25 min superheating, the ␣2 -Ti3 Al-phase was no longer detected, and the alloy is composed of only the ␥ and Ti2 AlN phases. The solubility of interstitial elements in the ␥ phase (L10 ) is known to be very limited. On the contrary the ␣2 -phase, having more loose-packed crystal lattice of D019 type, can dissolve
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typically two orders of magnitude more interstitials [12]. Moreover, nitrogen being in interstitial positions also stabilises the ␣2 -phase thermodynamically [13,14]. From these statements, it is clear that increasing the nitrogen content will at first lead to a rise of the ␣2 -phase fraction in the solidifying alloy, and also to a gradual competitive enrichment of this phase in nitrogen, since the ␥ phase can only dissolve a few hundreds of ppm. This “sponge effect” of ␣2 saturation with N finally leads to the formation of Ti2 AlN whose fraction rises with superheating time, and to the detriment of the ␣2 -phase. However, microstructure analysis has revealed the presence of one more newly formed phase. In order to identify it, some EBSD analyses were done in combination with EDX. These analyses have confirmed that the observed black precipitates (see again Fig. 2(b, d and f)) are indeed Ti2 AlN (Pearson symbol hP8, spacegroup P63 /mmc, No. 194), but they also revealed that the white particles are monoboride (Ti,Nb)B (Pearson symbol oC8, spacegroup Cmcm, No. 63). The reason why XRD did not reveal the boride phase is still not clear; probably it can be due to superimposing of some borides diffraction peaks with those of the ␥ phase. This would explain the measured volumetric ␥-fraction of 78.9% in the sample No. 3 whereas its value does not seem so high from viewing the microstructure. 3.3. Interaction dynamics The dynamics of reaction layer thickness evolution represents some interest as well. After 5 min of melt superheating this layer covers about 250 m at the bottom part of ingot, and still has a rather abrupt boundary (Fig. 2(a)). However, after 12 min it already exceeds 500 m, accompanied by an expanded transient diffuse area of boride/nitride defects precipitation (Fig. 2(c)). Finally, the interaction layer is not revealed at all after 25 min melt holding in BN crucible, since the whole sample (of 5 g mass) is filled with (Ti,Nb)B and Ti2 AlN micro-precipitates (Fig. 2(e and f)) and has changed drastically its phase composition. These results, in addition to the ones presented above, confirm that BN interacts quickly with molten TiAl-based alloys, and thus is not a suitable material as a mould for these highly reactive melts. 4. Summary and conclusions In this paper, we have investigated the chemical compatibility of Ti–46Al–8Nb melt with pyrolytic BN crucibles, in terms of cast alloy microstructure and contamination.
BN reacts with TiAl–Nb melts, leading to the formation of Ti2 AlN and (Ti,Nb)B needle-like micro-precipitates at the interface between alloy and crucible. The reaction layer formed by these particles propagates quickly towards the alloy, and affects the whole sample after 25 min superheating. Formation of both nitride and boride precipitates proceeds at the expense of selective destruction of the ␣2 -phase, being preceded by the “sponge effect” of accumulation of dissolved interstitial nitrogen and boron in a Ti3 Al sublattice. With some simplification, finally the scheme of alloy interaction with the BN crucible can be represented as follows: Ti3 Al + (Nb) + BN → Ti2 AlN + (Ti, Nb)B After 25 min of melt holding at 1670 ◦ C the amount of dissolved BN becomes so high that this reaction has apparently proceeded to completion and the former major Ti3 Al-phase is no longer detected in the 5 g sample of solidified alloy. Thus, it is shown experimentally that pBN is unsuitable material as a crucible/mould for highly reactive TiAl-based melts.
Acknowledgements This work was financially supported by EU Integrated Project IMPRESS under the contract No. NMP3-CT-2004-500635. A. Kartavykh acknowledges also the support of RFBR grant no. 0908-00844.
References [1] D.J. Jarvis, D. Voss, Mater. Sci. Eng. A 413–414 (2005) 583. [2] J.P. Kuang, R.A. Harding, J. Campbell, Mater. Sci. Technol. 16 (2000) 1007. [3] J.P. Kuang, R.A. Harding, J. Campbell, Int. J. Cast Met. Res. 13 (2001) 277. [4] A.V. Kartavykh, V.V. Tcherdyntsev, J. Zollinger, Mater. Chem. Phys. 116 (2009) 300. [5] URL: www.advceramics.com/products/pyrolytic bn/index.html. [6] URL: www.keramverband.de/brevier engl/3/4/3/3 4 3 2.htm. [7] T. Suzuki, A. Omino, J. Cryst. Growth 117 (1992) 80. [8] T. Katayama, T. Shibata, J. Cryst. Growth 24–25 (1974) 396. [9] X. Wu, D. Hu, M.H. Loretto, J. Mater. Sci. 39 (2000) 3935. [10] I. Egry, R. Brooks, D. Holland-Moritz, et al., Int. J. Thermophys. 28 (2007) 1026. [11] E.V. Shelekhov, T.A. Sviridova, Metal Sci. Heat Treatment 42 (2001) 309. [12] A. Menand, A. Huguet, A. Nerac-Partaix, Acta Mater. 44 (1996) 4729. [13] F. Perdrix, M.-F. Trichet, J.-L. Bonnentien, et al., Intermetallics 9 (2001) 147. [14] C.Y. Nam, D.M. Wee, P. Wang, K.S. Kumar, Intermetallics 10 (2002) 113.
Contents lists available at ScienceDirect
Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Materials science communication
TiAl–Nb melt interaction with pyrolytic boron nitride crucibles A.V. Kartavykh a,∗ , V.V. Tcherdyntsev b , J. Zollinger c,1 a
Institute of Chemical Problems for Microelectronics (IChPM), B. Tolmachevsky per. 5, 119017 Moscow, Russia Moscow State Institute of Steel and Alloys (Technological University), Leninsky pr. 4, 119049 Moscow, Russia c ACCESS e.V., Intzestrasse 5, 52072 Aachen, Germany b
a r t i c l e
i n f o
Article history: Received 31 March 2009 Received in revised form 10 September 2009 Accepted 18 September 2009 Keywords: Intermetallic compounds Solidification Ceramics Surface properties
a b s t r a c t This paper studies the chemical interaction on contact, and the related evolution of composition and microstructure of cast TiAl-based alloy, when using refractory oxygen-free pyrolytic boron nitride (pBN) crucibles/moulds instead of traditional oxide ceramic ones. We extend the research begun with AlN crucibles and reported in [4]. Three melting tests were performed in resistive SiC electro-furnace with Ti–46Al–8Nb (at%) alloy in 6N-purity pBN boat crucibles with fixed melt superheating times of 5, 12 and 25 min at 1670 ◦ C and subsequent quenching with high-purity Ar gas flow. As-cast samples were examined by optical microscopy/metallography, X-ray diffractometry, SEM-BSE, EDX and EBSD microanalyses, and compared to the initial material. Unlike the AlN it was shown that dissolved BN reacts with the melt, forming Ti2 AlN and (Ti,Nb)B micro-precipitates at the expense of selective destruction of the ␣2 -Ti3 Al major intermetallic phase in the solidified alloy: Ti3 Al + (Nb) + BN → Ti2 AlN + (Ti,Nb)B. The key features and temporal dynamics of this interaction are studied. © 2009 Elsevier B.V. All rights reserved.
1. Introduction TiAl-based alloys are good candidate materials for aero engine and gas-burning power-generating turbine blades [1]. One of the best processing routes for ␥-TiAl-based alloys is near-net shape investment casting in ceramic moulds in a high-purity inert gas atmosphere. The choice of mould material is a critical stage of the technological process. No refractories which are absolutely chemically inert against TiAl melts are known today [2,3]. The high solidification temperatures (1550–1600 ◦ C depending on alloying) and the melt reactivity both require more thermodynamically stable and chemically inert materials for mould manufacturing, instead of traditionally applied zirconia (ZrO2 ) and more advanced yttria (Y2 O3 -based) ceramics. One of the main reasons why the properties and structure of alloy being solidified are hard to reproduce is contamination with interstitial oxygen, which is difficult to control. The problem, mechanisms and consequences of uncontrolled oxygen contamination of TiAl intermetallics from crucibles/moulds at the melt processing have been observed elsewhere [4]. Thus, testing of alternative non-oxide (oxygenless) refractory materials may help in finding crucible materials which are better suited for processing the TiAl-based alloys. The results of such an experimental study, performed with the aluminium nitride
∗ Corresponding author. Tel.: +7 499 7889063; fax: +7 495 9538869. E-mail address: [email protected] (A.V. Kartavykh). 1 Present address: LETAM, Université Paul Verlaine de Metz, Ile du Saulcy, F-57012 Metz, Cedex, France. 0254-0584/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2009.09.021
crucibles, have previously been reported [4]. It was shown that AlN apparently holds promise as a mould material usable in metallurgical processing of TiAl intermetallics for the reduction of interstitial oxygen content. The nearest analogue to AlN in the series of refractory ceramics-like nitrides is boron nitride, having similar chemical and even higher thermal stability. This paper presents our investigations of the chemical compatibility of BN as the casting mould material for Ti–46Al–8Nb alloy. 2. Experimental procedures 2.1. Materials The pBN boat crucibles used in this work are serial commercial BoralloyTM products supplied from GE Advanced Ceramics, USA, having a volume of 2 cm3 , a mass of 1.5 g and wall thickness of 1.5 mm. They are manufactured by gas phase multilayer deposition, providing extra high-purity material of 6N grade [5]. The pyrolytic BN is stable up to 2500 ◦ C in argon; does not contain chemically bonded oxygen; it has zero open porosity, rather high thermal conductivity up to 60 W m−1 K−1 and linear thermal expansion coefficient of (2.2–4.4) × 10−6 K−1 promoting excellent resistance to thermal shock [5,6]. The pyrolytic BN is widely used in crucibles for processing of various intermetallics, alloys and semiconductor materials (see, for example [7,8]). Of particular importance is the fact that pBN maintains its resistance to corrosive attack by aluminum even at extreme temperature. Besides, boron (one of the constituents of the crucible material) is a routine grain refiner in as-cast TiAl-based alloys and is specifically used sometimes as a minor addition in the melt, at levels below 1 at.% [9]. Dissolved nitrogen (another constituent of crucible) does not have any detrimental effect on the microstructure of TiAl–Nb alloys, at least at contents less than 1800 wt ppm [4]. All the abovementioned prerequisites were taken into account in the choice of pBN as the possible alternative mould material for testing. The intermetallic was supplied through the IMPRESS project. Ti–46Al–8Nb (nominal composition, in at.%) alloy was prepared with induction melting and subsequent triple homogenizing vacuum arc remelting. Samples for tests have the cylindrical shape Ø 15 mm × 15 mm. A batch of 10 samples was delivered with the
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Table 1 Compatibility test conditions and XRD results for the initial Ti–46Al–8Nb material and as-cast samples. Sample number
Superheating time at 1670 ◦ C in pBN crucible (min)
Initial material No. 1 No. 2 No. 3
0 5 12 25
Phases revealed in the sample’s bulk and their volumetric fractions (%)
TiAl
Ti3 Al
Ti2 AlN
85.1 ± 0.2 84.8 ± 0.2 72.0 ± 0.5 78.9 ± 0.2
14.9 ± 0.2 15.2 ± 0.2 22.0 ± 0.5 0
0 0 6.0 ± 0.5 21.1 ± 0.2
Structural type of phase
tP2/1
hP8/3
hP8/4
Lattice periods (Å)
A = 2.841 ± 0.002 C = 4.084 ± 0.002
A = 5.780 ± 0.004 C = 4.654 ± 0.004
A = 3.003 ± 0.004 C = 13.628 ± 0.004
analysis certificate giving the real composition of Ti–46.35Al–7.87Nb (at.%). Our vacuum fusion analysis has established O, N, C and S interstitial impurities content of 700, 40, 230 and 25 wt ppm, respectively, in the alloy batch. 2.2. Experimental apparatus and conditions The experiments were performed in a special low-inertial electro-furnace, specified in detail in paper [4]. The facility consists of a vertical hermetic quartz flow-through reactor, connected to a gas-purifying line through the bottom flange. The reactor blow-through proceeds upward. Inside the chamber, a flat SiC/graphitecomposite resistive heater is installed, on which a boat crucible containing the test alloy sample may be directly placed. The sample temperature is measured by optical pyrometer with ±20 ◦ C accuracy and governed by heater power control. The small volume and mass of the test alloy charge (5 g) and crucible (1.5 g) both ensure the heat treatment of the sample occurs in near bulk-isothermal conditions. Each experiment was performed with the standard quartz tube-case being exchanged for a new one, to avoid any pollution accumulation in the reactor and to reduce the chemical background effect on material contamination level. During the experiment, technological argon was used. Ar passed through triple refining: drying in silica gel column following O2 and CO2 absorption in zeolite filters. The applied gas purification technique is standard for the semiconductor industry; its quality was controlled by measuring the dew point temperature, which did not exceed −70 ◦ C. Before each experiment, the charged furnace was evacuated down to ∼10−2 Torr pressure and then intensively blasted with Ar stream for 30 min in the cold state. Test cycles contained the following stages: heating of the crucible/sample pair up to 1670 ◦ C (i.e., 100 ◦ C in excess of the equilibrium liquidus of the test alloy [10]) at a rate of 120 ◦ C min−1 ; isothermal superheating of 5, 12 or 25 min duration, defining the fixed time of melt/crucible interaction; rapid cooling (quenching) at a rate of 15 ◦ C s−1 to preserve the interaction results. An Ar flow rate of 0.03 L s−1 at 120 kPa pressure was applied during the melting and superheating stages. Quenching was performed by upward blowing over the heater with a strong argon stream (five times increased Ar flow rate) while simultaneously switching off the power. The correspondence between the sample numbers and their superheating parameters is given in Table 1.
of the sample superheated for 5 min. The sample surface is mat black with apparently some BN scales on it—a first indication of an interaction between alloy and crucible. Fig. 1(b) shows that this mat black crust also exists on the free surface. The transverse vertical cut made reveals that the material has a metallic appearance, and also indicates good wetting between the melt and crucible from the meniscus shape of the sample free surface.
2.3. Analytical techniques Analytical techniques were applied to metallographic slices representing the vertical transverse cross-sections of samples (mini-ingots) (see Fig. 1). The 1-m diamond polishing of slices was done using Buehler equipment and consumables; no chemical etching was applied. Optical microscopy was performed using a POLYVAR Reichert–Jung microscope, with main the focus on identifying reaction layers in the vicinity of the former interface of the melt/crucible, and precipitates in the bulk of TiAl-based alloy. All reaction products revealed were subjected afterwards to electron-probe micro-analysis. The local phase composition was measured using scanning electron microscopy (SEM) and energy-dispersive X-ray spectrometry (EDX) techniques. For backscattered electron (BSE) imaging and for EDX a Gemini 1550 microscope was used with the INCA-Energy software. The crystal structure of precipitates was checked by electron backscatter diffraction (EBSD) using the crossbeam microscope 1540 XB and the software HKL Flamenco. For X-ray phase analysis (XRD), a DRON-3 diffractometer with CuK␣ radiation source was used, equipped with a specialized holder for study of bulk textured specimens with the angle of sample positioning at 30◦ . For averaging of the output diffraction, the rotation of each fixed TiAl–Nb slice was applied with respect to its normal. Spectra interpretation was carried out using software [11] allowing the analysis of relative phase content by the Rietveld method, with ±0.2% accuracy.
3. Results and discussion Fig. 1 gives a typical overview of as-processed sample. In Fig. 1(a) part of the pBN crucible is shown next to the crucible-adjacent side
Fig. 1. Photographs of (a) BN crucible and sample re-solidified after 5 min of the melt superheating (No. 1 in Table 1); (b) as-cut metallographic slice of the same sample.
A.V. Kartavykh et al. / Materials Chemistry and Physics 119 (2010) 347–350
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Fig. 2. Optical micrographs (on the left) and SEM-BSE microstructure images (on the right) at the interface between the reaction layer and Ti–46Al–8Nb alloy in samples processed at 1670 ◦ C in pBN crucible with (a and b) 5 min, (c and d) 12 min, and (e and f) 25 min superheating time. No definite boundary was established between the reaction layer and the alloy bulk in the latter sample whose whole volume is uniformly filled with needle-like precipitates.
3.1. Interaction layers and microstructure of cast alloy Microstructural evolution of the alloy as dependent on superheating time is depicted in Fig. 2. Pronounced interaction layers were found by optical microscopy in the samples superheated at relatively short durations, and clearly seen in Fig. 2(a and c). The transient zone between the reaction layer and the main bulk is shown in the samples held for 5 and 12 min in Fig. 2(b and d), using SEM-BSE mapping. The reaction layer can be observed at the bottom of these images and is composed of a black matrix with some white needles. On the upper part of these images, one observes the solidification microstructure of the Ti–46Al–8Nb alloy appearing in light grey contrast in the BSE mode. This microstructure is composed of the two intermetallic ordered phases ␣2 -Ti3 Al (D019 ) and ␥-TiAl (L10 ) that forms a lamellar texture. It can clearly be seen that black particles have formed ahead of the reaction layer front, whose volume fraction increases with superheating time. Moreover, in Fig. 2(d), some randomly distributed white needles appear in the whole microstructure. The microstructure of the sample after 25 min superheating is shown in Fig. 2(e and f) at equal magnifi-
cations. It can be seen from optical microscopy/metallography in Fig. 2(e) that the whole sample bulk is filled with short-needlelike precipitates. Its microstructure is composed of black and white particles in Fig. 2(f), and the residual alloy itself (appearing in grey BSE contrast in this image) has lost the lamellar structure and is composed now of a single phase. 3.2. Phase composition and evolution XRD analyses were performed on the bulk part of the as-cast samples as well as on the initial material. The results are shown in Table 1. The following events were observed with increasing superheating time: (i) after 12 min, the ␥/␣2 ratio decreased with respect to the initial material, i.e., the volume fraction of ␣2 increased. The as-cast alloy also contains up to 6% volumetric fraction of Ti2 AlN phase; (ii) after 25 min superheating, the ␣2 -Ti3 Al-phase was no longer detected, and the alloy is composed of only the ␥ and Ti2 AlN phases. The solubility of interstitial elements in the ␥ phase (L10 ) is known to be very limited. On the contrary the ␣2 -phase, having more loose-packed crystal lattice of D019 type, can dissolve
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typically two orders of magnitude more interstitials [12]. Moreover, nitrogen being in interstitial positions also stabilises the ␣2 -phase thermodynamically [13,14]. From these statements, it is clear that increasing the nitrogen content will at first lead to a rise of the ␣2 -phase fraction in the solidifying alloy, and also to a gradual competitive enrichment of this phase in nitrogen, since the ␥ phase can only dissolve a few hundreds of ppm. This “sponge effect” of ␣2 saturation with N finally leads to the formation of Ti2 AlN whose fraction rises with superheating time, and to the detriment of the ␣2 -phase. However, microstructure analysis has revealed the presence of one more newly formed phase. In order to identify it, some EBSD analyses were done in combination with EDX. These analyses have confirmed that the observed black precipitates (see again Fig. 2(b, d and f)) are indeed Ti2 AlN (Pearson symbol hP8, spacegroup P63 /mmc, No. 194), but they also revealed that the white particles are monoboride (Ti,Nb)B (Pearson symbol oC8, spacegroup Cmcm, No. 63). The reason why XRD did not reveal the boride phase is still not clear; probably it can be due to superimposing of some borides diffraction peaks with those of the ␥ phase. This would explain the measured volumetric ␥-fraction of 78.9% in the sample No. 3 whereas its value does not seem so high from viewing the microstructure. 3.3. Interaction dynamics The dynamics of reaction layer thickness evolution represents some interest as well. After 5 min of melt superheating this layer covers about 250 m at the bottom part of ingot, and still has a rather abrupt boundary (Fig. 2(a)). However, after 12 min it already exceeds 500 m, accompanied by an expanded transient diffuse area of boride/nitride defects precipitation (Fig. 2(c)). Finally, the interaction layer is not revealed at all after 25 min melt holding in BN crucible, since the whole sample (of 5 g mass) is filled with (Ti,Nb)B and Ti2 AlN micro-precipitates (Fig. 2(e and f)) and has changed drastically its phase composition. These results, in addition to the ones presented above, confirm that BN interacts quickly with molten TiAl-based alloys, and thus is not a suitable material as a mould for these highly reactive melts. 4. Summary and conclusions In this paper, we have investigated the chemical compatibility of Ti–46Al–8Nb melt with pyrolytic BN crucibles, in terms of cast alloy microstructure and contamination.
BN reacts with TiAl–Nb melts, leading to the formation of Ti2 AlN and (Ti,Nb)B needle-like micro-precipitates at the interface between alloy and crucible. The reaction layer formed by these particles propagates quickly towards the alloy, and affects the whole sample after 25 min superheating. Formation of both nitride and boride precipitates proceeds at the expense of selective destruction of the ␣2 -phase, being preceded by the “sponge effect” of accumulation of dissolved interstitial nitrogen and boron in a Ti3 Al sublattice. With some simplification, finally the scheme of alloy interaction with the BN crucible can be represented as follows: Ti3 Al + (Nb) + BN → Ti2 AlN + (Ti, Nb)B After 25 min of melt holding at 1670 ◦ C the amount of dissolved BN becomes so high that this reaction has apparently proceeded to completion and the former major Ti3 Al-phase is no longer detected in the 5 g sample of solidified alloy. Thus, it is shown experimentally that pBN is unsuitable material as a crucible/mould for highly reactive TiAl-based melts.
Acknowledgements This work was financially supported by EU Integrated Project IMPRESS under the contract No. NMP3-CT-2004-500635. A. Kartavykh acknowledges also the support of RFBR grant no. 0908-00844.
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