TiAl–Nb melt interaction with AlN refractory crucibles

Materials Chemistry and Physics 116 (2009) 300–304

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TiAl–Nb melt interaction with AlN refractory crucibles A.V. Kartavykh a,∗ , V.V. Tcherdyntsev b , J. Zollinger c 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 29 October 2008 Received in revised form 25 January 2009 Accepted 26 March 2009 Keywords: Intermetallic compounds Solidification Ceramics Surface properties

a b s t r a c t The paper considers contamination of cast TiAl-based alloy, related microstructure evolution and chemical contact interaction when using refractory aluminium nitride crucibles/moulds as an alternative to the traditional Al2 O3 -, ZrO2 - and Y2 O3 -based oxide ceramics. A series of melting tests has been performed in resistive SiC electro-furnace with Ti–46Al–8Nb (at%) alloy in 99.99% purity AlN boat crucibles with fixed melt superheating times 5, 12 and 25 min at 1670 ◦ C and consequent quenching with high-purity Ar gas flow. As-cast samples were examined by X-ray diffractometry, SEM, SIMS, EDX, EBSD and vacuum fusion analysis of O, N, C and S interstitials content with respect to the initial material. The key features of TiAl–Nb melt interaction with AlN ceramics are revealed. As a result of slow thermal dissociation AlN → Al + N, and the reaction of nascent nitrogen with the melt, a solid continuous TiN-based reaction layer is formed up to 6.4 ␮m in thickness, together with an enriched Al liquid film between it and the crucible wall. It causes perfect wetting of crucible with the melt and easy removal of the solidified sample. The partial suitability of AlN crucible is restricted by 12 min of the melt superheating from the point of view of invariable (␣2 + ␥)-microstructure and reasonable contamination of the as-cast samples. © 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 manufacturing [1–3]. One of the perspective processing routes of ␥-TiAl-based alloys is near-net shape investment casting in ceramic mould under high-purity inert gas atmosphere. The choice of mould material is a critical stage of the technological process. No absolutely chemically inert refractories against TiAl-melts are known today [4,5]. The high solidification temperatures (1550–1600 ◦ C depending on alloying), and the melt reactivity both request 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 of properties and structure irreproducibility of alloy being solidified is hardly controlled contamination with interstitial oxygen. Oxygen content is mainly attributed to the thermo-chemical erosion/dissolution of ceramic; its increasing in Ti–45Al alloy to 1.5 at% changes the primary solidification phase from ␤ to ␣ [6]. When melting TiAl alloys in oxide crucibles, an oxygen-containing micro-precipitates are always present in the bulk [7], and some dispersed mechanical inclusions of eroded ceramics can also be found [8,9]. Their size

∗ Corresponding author: Tel.: +7 499 7889063; fax: +7 495 9538869. E-mail address: [email protected] (A.V. Kartavykh). 0254-0584/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2009.03.032

and amount both depend on the melt superheating temperature and time. These particles deteriorate sufficiently the ductility and fracture toughness of material, increase its fragility and reduce as a result the quality and runtime of industrial item produced. The elements bounding into oxide precipitates apart from Al and Ti may be other diffusing from crucible wall metals (Y, Zr, Ca, Mg, Si, etc. depending on the kind of binder and the composition of ceramics applied). Testing of alternative, non-oxide refractory materials may contribute to find better suited crucible materials for processing the TiAl-based alloys. The paper presented investigates the compatibility of aluminium nitride as the casting mould material for Ti–46Al–8Nb alloy. 2. Experimental 2.1. Materials High-purity AlN does not contain chemically bounded oxygen; it is inert against Al melt; it does not contain a foreign metal relatively to TiAl-based melts; it has a unique for ceramics heat conductivity of 180–220 W m−1 K−1 and linear expansion coefficient of (4.5–5.6) × 10−6 K−1 promoting good resistance to thermal shock [10]. AlN is stable up to 1800 ◦ C in argon atmosphere at normal pressure; it decomposes completely in Ar at 2500 ◦ C [11]. The main decomposition reaction of AlN is thermal dissociation into components [10–12]. However, AlN melts without decomposition under high pressure of nitrogen [11] that allows crucible manufacturing without any binders. The AlN boat crucibles were manufactured at the State Research Rare Metal Institute (GIREDMET), Moscow, from 99.99% purity AlN powder by means of hightemperature sintering/melting in N2 atmosphere at 2100 ◦ C and 3 MPa pressure. The intermetallic was supplied through the IMPRESS project. Ti–46Al–8Nb (nominal composition, in at%) advanced alloy was prepared with induction

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Fig. 1. The heating unit of the test furnace ready for experiment: 1—targeted points of optical pyrometer for heater and sample (melt); 2—planar resistive heater made from SiC–graphite composite; 3—boat crucible with a sample; 4—water-cooled power suppliers; 5—bottom flange of reactor (stainless steel) with case-hermetizing rubber ring; 6—inlet pipe for Ar flow into reactor and quench of heater by gas stream; 7—furnace base.

melting and subsequent triple homogenizing vacuum arc remelting. 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. Apparatus and conditions For experiments a special low-inertial electro-furnace has been used. The facility represents itself the vertical hermetic quartz flow-through reactor of 800 mm length and an inner diameter of 85 mm, connected to gas-purifying line through the bottom flange. The reactor flow-through proceeds upward. Inside the chamber, on 1 ⁄3 of its height the flat resistive heater is installed made from SiC/graphite composite, on which a boat crucible may be directly placed with the testing alloy sample. The general view of charged furnace heat unit is shown in Fig. 1. The heater connection to D.C. 6 V stabilized power source is performed through the copper hollow-rod supplies with inner water cooling. The heater joins to power supplies by means of threaded connections, intermediate bush inserts and bolts also machined from multilayer composite material “SiC/graphite”. The temperature of sample is measured by optical pyrometer with ±20 ◦ C accuracy and governed by heater power control. The small volume and mass of testing alloy charge (5 g) and crucible (2 g) both ensure the heat treatment of a sample in near bulk-isothermal conditions. Each experiment was performed with the standard quartz tube-case being exchange to a new one for avoiding any pollution accumulation in the reactor and for reducing the chemical background effect on material contamination level. The technological argon passed through triple refining, which is standard for semiconductor industry, i.e. drying in silica gel column following O2 and CO2 absorption in zeolite filters. The quality of gas purification 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. with 100 ◦ C over the equilibrium liquidus of testing alloy [13]) with

Fig. 2. (a) As-cut vertical sections of the cast samples: the upper meniscus shape exhibits the perfect wetting of AlN boat crucible with TiAl–Nb melt; (b) the exterior of sample No. 3 in top and bottom view. The rest of the continuous “gold” reaction layer is clearly seen on the crucible-adjacent surface (right). The fragmentation and partial exfoliation of this layer is caused by shrinkage of alloy (whose solid/melt density difference is 3.98/3.85 g cm−3 [13] near the melting point), and by final quench-induced thermal stress. 120 ◦ C min−1 rate; isothermal superheating of 5, 12 or 25 min duration, defining the fixed time of melt/crucible interaction; rapid cooling (quenching) with 15 ◦ C s−1 rate to conserve 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 performed by upward blowing 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. 2.3. Analytic techniques Analytic techniques were mainly applied on metallographic slices representing the vertical transverse cross-sections of samples (mini-ingots) (see Fig. 2a). The 1-␮m diamond polishing of slices was done using the Buehler equipment; no chemical etching was applied. The elemental profile analysis of reaction layers and evaluation of their thickness was performed on as-cast ingots by secondary ions mass-spectrometry (SIMS) with direct on-depth ion etching using IMS-4F CAMECA device. The local phase composition was measured using scanning electron microscopy (SEM) and the energy-dispersive X-ray spectrometry (EDX) techniques. For backscattered electron (BSE) imaging and for EDX the Gemini 1550 microscope was used with the INCA-Energy software. The crystal structure of precipitates was checked by electron backscatter diffraction (EBSD) using the cross-beam microscope 1540 XB and the software HKL Flamenco. For X-ray phase analysis (XRD) DRON-3 diffractometer was applied with CuK␣ radiation source, equipped with the specialized holder for bulk textured specimens study with the slope angle of sample

Table 1 The reaction contact layer thickness and phase composition of Ti–46Al–8Nb alloy after the compatibility tests. Sample number

Superheating time at 1670 ◦ C in AlN crucible (min)

Thickness of TiN reaction layer (␮m)

Initial material No. 1 No. 2 No. 3

0 5 12 25

n/a 1.1–2.0 3.1 6.4

Volumetric fraction of phases revealed with XRD analysis (%) ␥-TiAl

␣2 -Ti3 Al

85.1 85.2 85.0 80.9

14.9 14.8 15.0 19.1

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2–4 at%, correspondingly, along with the trace amount (0.5–1 at%) of nitrogen being also detected. From all the presented above, the possible mechanism of reaction of AlN crucible with TiAl–Nb alloy is as follows: 1) At 1670 ◦ C the slow thermal dissociation of aluminium nitride takes place: AlN (solid) → Al (melt) + 12 N2

(1)

2) The nascent atomic nitrogen is highly reactive and bounds immediately with titanium at the crucible/melt interface, forming a dense solid reaction layer (coating) of titanium-based nitride: (Ti, Al, Nb) (melt) + 12 N2 ↔ TiN (solid) + (Ti, Al, Nb, N) (melt) (2) The rate of decomposition (1) is temperature dependent; it is constant under conditions of isothermal exposure. This way, at the boundary between the crucible and the growing nitride layer, an excess of atomic nitrogen is permanently created, that shifts the reaction (2) to the right side according to Le Chatelier principle. The proceeding of reaction (2) exhibits the higher thermodynamical stability of TiN in comparison with AlN at 1670 ◦ C. Therefore, the summary chemical process may simply be considered like the slow reaction of Al replacement by Ti: AlN (solid) + Ti (melt) → TiN (solid) + Al (melt)

Fig. 3. SIMS profiles from the reaction layer towards the bulk in sample No. 3, showing the qualitative elements distribution with the depth; the layer thickness is 6.4 ␮m. positioning of 30 ◦ . For averaging of output diffraction the rotation of each fixed TiAl–Nb slice was applied with respect to its normal. Spectra interpretation has been done with the software [14] allowing the analysis of relative phase content by the Rietveld method with ±0.2% accuracy. The analysis of overall oxygen, nitrogen, carbon and sulphur contents in sample’s bulk was performed by vacuum fusion technique with ±20 % accuracy using LECO PO-316 device. For analysis the cubic blocks were used with 4 mm ridge being cut from the as-cast and initial samples.

(3)

Let us note once more that the layers resulted from TiAl–Nb/AlN interaction are quite not ordinary and represent continuous, rather perfect film of several microns thickness, deposed on the crucible-adjacent ingot side. A deeply extended penetration zone of damaged structure has not been formed in an ingot contact area, in contrast with the cases of oxide crucibles application for melting of TiAl-based alloys, known from the literature. The thickness of interfacial interaction layer has the order of several tens or even hundreds microns [4,5,9] when using an oxide crucibles in similar tests. Reaction (3) of TiN coating formation at the crucible/melt contact interface is the key feature of TiAl-based melts solidification in AlN crucibles. It leads to perfect wetting ability of the crucible with the melt (see Fig. 2a), avoiding the appearance of cast surface

3. Results and discussion The free upper surface of all as-cast mini-ingots has a metallic aspect, whereas the bottom part has a golden glance (Fig. 2b). According to EDX and SIMS data, the main substance of this “gold” deposed layer is TiN. The overall thickness of the coating grown at the longest superheating time was measured as 6.4 ␮m (Fig. 3). One can distinguish 2 distinct areas within the layer: from the outer side, the profile displays depletion in Al and Nb over about 3 ␮m, in the area that tally with TiN. The others 3 ␮m are in fact a transient diffusion area. In the alloy the amount of nitrogen falls below the SIMS detection limit within two more microns. The dynamics of reaction layer formation is represented in Table 1. The dependence of its thickness on superheating time is practically linear. After removing the samples, the extended fragments of metallization were observed at the inner surface of a crucible after each experiment (see Fig. 4). EDX analysis confirmed that they are strongly enriched with Al with respect to the tested alloy composition. The averaged fraction of aluminium in this metallic film is consisted 80–85 at% while the contents of Ti and Nb are reduced down to 12–17 and

Fig. 4. Fragments of Al-based metallic film found on the inner surface of AlN crucible, resulted from the slow high-temperature dissociation of AlN.

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Fig. 5. Microstructure overview (SEM–BSE) at the bottom area of sample No. 3.

defects (pores) (see Fig. 2b), and at the same time it ensures the easy withdrawal of ingot solidified. No other phases apart from the main intermetallic ones, i.e. from ␣2 -Ti3 Al (D019 ) and ␥-TiAl (L10 ) were revealed in the bulk of samples with XRD. However, the results shown in Table 1 for sample No. 2 and No. 3 point out an increase of the volume fraction of ␣2 phase from 15 to 19% at the expense of the gamma phase, which can be an indicator of nitrogen (or others ␣2 -phase stabilizing interstitial elements) pickup that is in good agreement with the observations of [15]. It is well-known that the elements N, O, C, S are interstitials in Ti3 Al sub-lattice, and they stabilise thermodynamically this phase in TiAl. The cumulative level of interstitials content for sample No. 2 is established as 3620 wt ppm (where corresponding O, N, C and S contents are 1300, 1800, 490 and 30 wt ppm), and for sample No. 3 is 5640 wt ppm (where O, N, C and S consist 1800, 3000, 800 and 40 wt ppm, respectively). So, one can propose that the latter general contamination becomes sufficient for a noticeable shift of ␥/␣2 -phase balance in the as-cast material with respect to the initial one. The lamellar solidification microstructure composing of two ordered phases Ti3 Al and TiAl is clearly visible in BSE images for all samples. No evident contamination was observed after 5 and 12 min superheating, whereas a few black particles are appeared in 100␮m near-bottom zone of re-solidified ingot after 25 min processing (Fig. 5). They are identified as TiN (Pearson symbol cF8, spacegroup Fm-3m, No. 225) with EBSD, and EDX analyses have shown neither Al nor Nb content in these particles. It is important that no mechanical inclusions of crucible-eroded ceramics were found in material bulk unlike to the cases of oxide compounds use [8,9]. Fig. 6 shows the increasing of N and O pickup by alloy with superheating time in the present work, along with the literature data for samples prepared in oxide crucibles. Only the additional contents of interstitials incorporated into alloy in result of the test melting are represented here. From this figure, one can conclude that the nitrogen pickup in the melt is following a fairly parabolic asymptotic, which can be attributed to the barrier effect of the dense TiN-layer between the crucible and the melt. One also observes a maximum level (1100 wt ppm) of capture of residual oxygen from the crucible pores (or, most probably, from the porous SiC/graphite material of the heater) in the sample No. 3. This content is lower than that reported in similar experiments for pure CaO [4], pure Al2 O3 [5] and Y2 O3 -coated zirconia [9] crucibles, and some higher than that for dense yttria use [5]. Note however, that the values reported from [4,5,9] were obtained at 1550 ◦ C that is 120 ◦ C lower than in the present work. For sample No. 2 the ␥/␣2 -phase balance

303

Fig. 6. The major interstitials pickup by alloy vs. the superheating time for AlN crucible use in comparison with the literature data for oxide crucibles (PW: present work).

is still preserved, despite the pickup of oxygen and nitrogen consist 600 and 1760 wt ppm, respectively. At this contamination level nitrogen is still in solid solution form and does not have any detrimental effect on the mechanical properties of ␥-TiAl-based alloys [15,16]. 4. Summary and conclusions The paper has investigated the physical–chemical compatibility of Ti–46Al–8Nb melt with AlN refractory crucible ceramics. The permanent simultaneous formation of neighbouring aluminium layer (from the side of crucible) and TiN-based solid film (from the side of melt) at the crucible/melt interface is the key event at the temperature investigated (1670 ◦ C): (i) it leads to good wetting ability of crucible with the melt, avoiding the appearance of cast surface defects (pores); (ii) it ensures the easy withdrawal of the solidified ingot. This nitride reaction layer has the thickness of order of few microns and can be easily removed by sand blasting or mechanical machining at the finish treatment of cast items. The main advantage of AlN crucible use in comparison with an oxide ceramics is twofold: (i) no deeply extended diffusion penetration zone of damaged structure formed in an ingot contact area; (ii) no mechanical inclusions of crucible-eroded ceramics found in the alloy matrix. The metallurgical quality of the samples superheated in AlN crucible during 5 and 12 min and re-solidified seems rather satisfactory in terms of invariable microstructure and general contamination. Acknowledgements This work was financially supported by EU Integrated Project IMPRESS “Intermetallic Materials Processing in Relation to Earth and Space Solidification” under the contract No. NMP3-CT2004–500635. A. Kartavykh acknowledges also the support of RFBR grant No. 09-08-00844 for partial test equipment funding. References [1] [2] [3] [4] [5] [6] [7]

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