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Experimental study of the Fe–Ni–Ti system
Intermetallics 18 (2010) 374–384
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Experimental study of the Fe–Ni–Ti system L.I. Duarte a, b, *, U.E. Klotz a,1, C. Leinenbach a, M. Palm c, F. Stein c, J.F. Lo¨ffler b a
¨ berlandstrasse 129, CH-8600 Du ¨ bendorf, Switzerland Empa, Swiss Federal Laboratories for Materials Testing and Research, Laboratory for Joining and Interface Technology, U Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, CH-8093 Zurich, Switzerland c ¨ r Eisenforschung GmbH, Max-Planck-Str. 1, D-40237 Du ¨ sseldorf, Germany Max-Planck-Institut fu b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 8 August 2008 Received in revised form 12 August 2009 Accepted 13 August 2009 Available online 3 September 2009
In this investigation phase relations in the Fe–Ni–Ti system were studied and two isothermal sections at 800 C and 1000 C as well as a revised liquidus projection were established. Microstructural characterisation of the as-cast alloys and of samples equilibrated at 800 and 1000 C was performed by scanning electron microscopy (SEM), chemical compositions of the phases were analysed by electron probe microanalysis (EPMA), and liquidus temperatures were examined by differential thermal analysis (DTA). The experimental results clarify some uncertainties concerning the melting behaviour and the solid-state phase equilibria between the phases (Ni,Fe)Ti2, (Fe,Ni)Ti and b-Ti. The present data also confirm that the solid solubility of Ti in g-(Fe,Ni) varies in dependence on the Fe:Ni ratio and decreases with decreasing temperature. The liquidus projection as well as the reaction scheme in the Ti-lean part are modified because two ternaries eutectic E1: L 4 g-(Fe,Ni) þ Fe2Ti þ Ni3Ti and E2: L 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti are found at 1108 and 1099 C, respectively. Ó 2009 Elsevier Ltd. All rights reserved.
Keywords: A. Ternary alloy systems A. Laves phases B. Phase diagrams F. Microprobe
1. Introduction Fe–Ni–Ti is one of the ternary subsystems of the Al–Fe–Ni–Ti quaternary system, which is relevant for many different materials such as steels, Ni-based superalloys, alloyed aluminides, Ti-based corrosion-resistant alloys, shape-memory alloys, and amorphous materials. The Fe–Ni–Ti system has been studied extensively by a large number of investigators and was reviewed by Gupta [1] and later revised by the same author in Ref. [2]. More recently, this system was reviewed by Cacciamani et al. [3] and by Ghosh [4]. The liquidus projection and two isothermal sections at 900 C and 1000 C are known from literature, as discussed below. Gupta [1] proposed a schematic liquidus projection based mainly on the work of Vogel and Wallbaum [5], Speich [6], and Van Loo et al. [7]. The first comprehensive study of phase equilibria from the Fe–Ni–Ni3Ti– Fe2Ti part of this system was carried out in Ref. [5]. Thermal analysis and microstructural investigations were performed on several
* Corresponding author. Empa, Swiss Federal Laboratories for Materials Testing ¨ berlandstrasse 129, and Research, Laboratory for Joining and Interface Technology, U CH-8600 Du¨bendorf, Switzerland. Tel.: þ41 44 823 42 28; fax: þ41 44 823 40 39. E-mail address: [email protected] (L.I. Duarte). 1 Present address: FEM, Research Institute for Precious Metals and Metals Chemistry, Department of Physical Metallurgy, Katharinenstrasse 17, D-73525 Schwa¨bisch Gmu¨nd, Germany. 0966-9795/$ – see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2009.08.008
alloys and the results were reported in vertical sections (isopleths). From these results it was possible to draw a partial liquidus projection and several vertical sections. Later, Dudkina and Kornilov [8] established a vertical section TiFe–TiNi using twenty-one ternary alloys. They determined the solidus along this section and found complete solubility of Fe and Ni in the ordered phase (Fe,Ni)Ti. Alisova et al. [9] provided three vertical sections in the Tirich corner (>50 at.% Ti) for the Fe:Ni atomic ratios 3:1, 1:1, and 1:3 and a partial reaction scheme of the invariant equilibria in the Ti– TiNi–TiFe region. According to this work a four-phase transitiontype (U) reaction L þ (Fe,Ni)Ti 4 b-Ti þ NiTi2 takes place at 960 C and very close to the binary eutectic L 4 b-Ti þ NiTi2. The vertical sections provided in [9] are in disagreement with the results reported by Efimenko et al. [10]. In that work, diffusion couples were used which were produced by melting a mixture of Ni and Ti powders on an iron substrate and annealing at 1100, 1200 and 1250 C for 5–120 min. In the solidified melts, the two phases NiTi2 and (Fe,Ni)Ti were always observed, although according to Alisova et al. [9] an equilibrium between b-Ti and (Fe,Ni)Ti would have been expected. In the recent assessment by Gupta [2] the three vertical sections provided by Alisova et al. [9] were taken into consideration. However, the reaction L þ (Fe,Ni)Ti 4 b-Ti þ NiTi2 was shifted towards a much lower Fe content than in a previous assessment. Nevertheless, it was not stressed that the results from [9] disagree with the isothermal section at 900 C provided by van Loo et al. [7], who showed NiTi2 to have a high solubility for Fe at
L.I. Duarte et al. / Intermetallics 18 (2010) 374–384
900 C. Ref. [7] provides the complete isothermal section at 900 C, showing that no ternary compound exists at this temperature and that all the binary compounds, i.e. NiTi2, NiTi, FeTi, Fe2Ti and Ni3Ti, possess a significant solubility for the third element. Abramycheva et al. [11] determined an isothermal section at 1000 C based on experimental results from several diffusion couples. In addition, a number of data have been reported for a Ti content of less than 50 at.% [6,12–14]. Speich [6] and Fournelle [12] studied the cellular precipitation of Ni3Ti in g-(Fe,Ni) in alloys with 28.7 at.% Ni and 7.0 at.% Ti. Speich [6] found a solvus at 982 C and a lower solubility limit of Ti in the g-(Fe,Ni) phase compared to Ref. [11]. Abraham [13,15] also found a lower solubility limit of Ti in the g-(Fe,Ni) phase determined at 1024 C in an alloy containing 27 at.% Ni. The scattering of the values for the solubility of Ti in the g-(Fe,Ni) phase was previously discussed in Ref. [1] and compared with the work done by Drake [16], where it was found that the maximum solubility of Ti in g-(Fe,Ni) at 1027 C was higher than 12–13 at.%. Consequently, the boundary of the g-(Fe,Ni) phase field was redrawn taking into consideration the work of Ref. [11]. It is noted that phase equilibria between g-(Fe,Ni), a-Fe and Fe2Ti (and L) at 1100 and 1200 C were recently determined by Sugiura et al. [17]. The isothermal section at 1000 C appears to be similar to that at 900 C for Ti content lower than 50 at.%. However, a number of differences exist in the Ti-rich part: from the binary Ni–Ti system the liquid phase extends into the ternary system at 1000 C. Also, according to [10], the (Ni,Fe)Ti2 phase, found in connection with very high Fe content, shows a very small homogeneity range at 1000 C. In addition, the nature and the position of the invariant reaction between the melt, (Ni,Fe)Ti2 and (Fe,Ni)Ti remains unclear. According to [11] the stability of (Ni,Fe)Ti2 at 1000 C is limited to an area close to Fe–Ti. An equilibrium between the liquid phase and (Ni,Fe)Ti2 was not observed. Instead, a three-phase equilibrium between b-Ti, (Ni,Fe)Ti2 and (Fe,Ni)Ti appears to exist at this temperature. In consequence, the field of primary solidification of (Ni,Fe)Ti2 must be substantially extended and the invariant reaction is not close to the Ni–Ti side but to the Fe–Ti side. If this is the case, two reactions are possible: an invariant peritectic (P) reaction with L þ b-Ti þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 or a ‘‘quasi-peritectic’’ four-phase transition-type (U) reaction with L þ (Fe,Ni)Ti 4 b-Ti þ (Ni,Fe)Ti2 associated with a maximum in the monovariant liquidus line connecting it with the binary peritectic L þ NiTi 4 NiTi2, as suggested in Ref. [3]. Riani et al. [18] chose the second option to explain their recent experimental results, obtained from DTA measurements and isothermal annealing at 900 C. The objective of this study is to determine the isothermal sections at 800 C and 1000 C experimentally and to clarify the uncertainties in the Ti-rich corner of this system. The determination of phase equilibria and solidification sequence at 66.7 at.% Ti is also an objective, to find out how far NiTi2 is stabilized by Fe. The solidus and liquidus temperatures of the Ti-rich alloys will provide a basis for a more complete thermodynamic description of the Fe–Ni–Ti system. 2. Experimental Pure elements of Fe (99.97%), Ni (99.98%) and Ti (99.99%) were used to produce the alloys given in Table 1. Alloys of about 4 g, most of them in the Ti-rich corner, were prepared by arc-melting under a purified argon atmosphere (99.999%) using a non-consumable tungsten electrode. All alloys were melted five times and inverted after three meltings to ensure homogeneity. In order to have a clean atmosphere during the melting, a Ti alloy was melted first as an oxygen getter. In addition, an oxygen cartridge in the argon line was used. No chemical analysis of the alloys was conducted since the weight losses during melting were less than 0.2 mass% in all cases. Afterwards, the samples were cut into 4 pieces for various
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Table 1 Nominal compositions of the Fe–Ni–Ti alloys produced by arc-melting, temperatures and annealing times of the heat treatments, and onset temperatures of melting and liquidus temperatures measured during heating of as-cast alloys by DTA. Nominal composition, (at.%) Alloy no.
Ni
01 02 03
2.7 74.1 23.2 800 19.5 69.4 11.1 1000 2.0 66.6 31.4 1000 800 6.5 66.6 26.9 800 10.3 66.5 23.2 1000 800 15.8 66.5 17.7 1000 800 4.6 58.4 37.0 – 44.4 16.6 39.0 1000 800 55.4 15.0 29.6 1000 800 65.3 14.0 20.7 1000 800 29.6 20.4 50.0 1000 800 9.4 22.8 67.8 1000 800 2.8 12.0 85.2 1000 800 0.9 13.9 85.2 1000 800 14.8 80.6 4.6 1000 800 19.2 61.7 19.1 1000 800 6.7 45.0 48.3 1000 800 27.2 40.0 32.8 1000 800 43.5 34.0 22.5 1000 800 56.9 36.9 6.2 1000 800 5.0 66.6 28.4 – 10.0 66.7 23.3 1000 9.0 72.2 18.8 1000 15.0 66.7 18.3 1000 25.1 66.6 8.3 1000
04 05 06 07 08 09 10 11 12 13 14 15 16 18 19 20 21 32 33 34 35 36
Ti
Fe
Temperature ( C)
DTA temperatures from heating ( C) Annealing time Onset Liquidus (h) temperature 1119 10 min 1167 1119 1119 1167 1119 375 375 – 1167 1119 1167 1119 1167 1119 1167 1119 1167 1119 1167 1119 1167 1119 10 min 375 375 375 1167 1100 1167 1100 1167 1100 1167 1100 – 702 702 702 702
1052 1128 988 1083 1062 1128 1028 1153 1020 1143 1024 1136 1044 1109 1252 1211 1267 1108 1179 1199 1270 1267 – 1282 – 975
–
1029 1160 –
–
1204 – 1099 1150 1113 1201 1040 1023 1022 1017 1018
1152 1158 1098 1159 1112
experiments. One piece was used to investigate the as-cast microstructure; another was used for differential thermal analysis (DTA) experiments; and the remaining two were used to investigate the microstructure after annealing experiments at 800 and 1000 C. For the heat treatments the samples were placed in evacuated quartz tubes re-filled with argon. Before the encapsulation the samples were cleaned in acetone, dried and wrapped in tantalum foil to avoid reactions between samples and the quartz tube. A titanium getter was placed in the quartz tube to absorb oxygen. The annealing experiments were performed at 800 C and 1000 C for 375 to 1200 h (c.f. Table 1) in an electric resistance furnace with a temperature accuracy of 3 C. After annealing, the samples were quenched in salt water to retain the equilibrium microstructures. In order to study the equilibrium between the liquid phase and NiTi2 at 1000 C a diffusion couple was prepared from an arcmelted alloy with the chemical composition of NiTi2 and a crucible of pure Fe (99.8%). The Fe crucible had a cylindrical shape of 20 mm diameter and 20 mm in height, with a hole of 16 mm in diameter and 17 mm in length. A NiTi2 sample of 16 mm in diameter and 5 mm in height was inserted into the Fe crucible. The diffusion couple was annealed for 30 min at 1010 C in a vacuum furnace (Torvac 1218H-PQ) under a pressure of <0.001 Pa.
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The heating rate was 10 C/min and the cooling rates were 50 C/min and 15 C/min from 1010 C to 800 C and from 800 C to room temperature, respectively. Higher cooling rates could not be achieved due to furnace limitations. After the annealing experiments, the samples were cut transversally by electro-discharge machining (EDM) and prepared by standard metallographic methods. The as-cast and equilibrated alloys were then examined by optical microscopy and scanning electron microscopy (SEM). The determination of the compositions of the coexisting phases of each sample was performed by electron probe microanalysis (EPMA) using a JEOL JXA8800 microanalyser with an acceleration voltage of 20 kV and a probe current of 2 108 A. For the measurement of the eutectic areas the beam was widened to 3–5 mm. Pure element standards provided by JEOL were used for calibration. The measurements have a relative accuracy of about 1%. If standard deviations given in Tables 2 and 3 are higher than this value or are missing only a few points could be measured, indicating that the phase was either too small to measure or that only a few particles could be detected. DTA experiments were performed in a Setaram SETSYS-18 and in a Netzsch STA 409 CD. Alumina crucibles coated with yttrium oxide were used in order to avoid reactions between the alloy and the crucible. The DTA cells were calibrated using the melting temperatures of the pure elements Al, Ag and Au. Cylindrical samples of 3 mm in diameter and a height of 5 mm, which were cut by EDM from equilibrated samples, were used in the DTA experiments. Before each experiment, the DTA cell was evacuated three times and re-filled with high-purity argon. Measurements were performed at a scanning rate of 10 C/min for heating and cooling under a continuous flow of argon. The evaluation of the DTA curves showed that the samples still exhibited some reactions with the crucibles due to the high reactivity of Ti in the liquid. Therefore only the data from the first heating curve were taken into account. 3. Results The overview in Table 1 shows the compositions of the alloys prepared by arc-melting, the temperatures and durations of the heat treatments, the onset temperatures of melting and the liquidus temperatures obtained from the DTA experiments. 3.1. As-cast alloys
Table 2 Compositions of coexisting phases measured by EPMA for alloys annealed at 1000 C and for the diffusion couple. Alloy no.
Phase
Individual alloys b-Ti 03 (Fe,Ni)Ti b-Ti 02 (Ni,Fe)Ti2 Liquid b-Ti 05 (Ni,Fe)Ti2 (Fe,Ni)Ti 06 (Ni,Fe)Ti2 (Fe,Ni)Ti 08 Ni3Ti g-(Fe,Ni) 09 Ni3Ti g-(Fe,Ni) 10 Ni3Ti g-(Fe,Ni) 11 Fe2Ti Ni3Ti g-(Fe,Ni) 12 Fe2Ti g-(Fe,Ni) 13 Fe2Ti a-Fe 14 Fe2Ti a-Fe b-Ti 15 (Ni, Fe)Ti2 Liquid 16 (Ni,Fe)Ti2 (Fe,Ni)Ti 18 (Fe,Ni)Ti Fe2Ti 19 (Fe,Ni)Ti Fe2Ti 20 (Fe,Ni)Ti Fe2Ti Ni3Ti 21 Ni3Ti (Fe,Ni)Ti 33 (Ni,Fe)Ti2 (Fe,Ni)Ti b-Ti 34 (Ni,Fe)Ti2 35 (Ni,Fe)Ti2 36 (Ni,Fe)Ti2 (Fe,Ni)Ti
Ni (at.%)
Ti (at.%)
Fe (at.%)
1.2 0.1 2.9 0.1 8.6 0.7 19.9 0.3 n.d. 3.8 0.1 10.2 0.1 9.4 0.1 15.1 0.1 18.1 0.1 59.7 0.9 30.7 0.7 66.0 0.5 40.8 0.4 71.1 0.4 59.7 0.3 18.9 0.3 56.8 0.7 28.4 0.4 8.7 0.1 11.2 0.1 3.0 0.1 2.7 0.1 0.9 0.1 0.7 0.1 8.8 0.5 24.5 0.5 n.d. 17.3 0.1 22.4 0.1 8.6 0.1 3.3 0.7 32.4 0.3 19.2 0.2 40.9 0.1 26.0 0.1 65.4 0.3 70.2 0.1 46.2 0.1 10.8 0.1 10.2 0.1 4.1 0.1 11.8 0.4 15.0 0.1 23.9 0.3 33.8 0.3
78.1 0.4 53.1 0.4 85.3 0.4 68.2 0.4 n.d. 77.3 0.1 67.2 0.3 52.4 0.1 67.4 0.3 52.4 0.1 23.8 0.5 10.3 0.4 23.1 0.4 7.1 0.2 22.0 0.4 6.3 0.6 28.4 0.6 24.3 0.1 12.0 0.2 28.2 0.3 6.9 0.4 28.5 0.2 5.7 0.1 28.8 0.1 5.3 0.2 86.9 0.3 69.0 0.4 n.d. 67.5 0.1 52.1 0.1 50.1 0.2 35.8 1.8 46.4 0.3 34.8 0.2 43.0 0.2 34.1 0.1 26.4 0.2 26.3 0.1 44.8 0.1 67.7 0.2 53.1 0.1 79.9 0.2 67.9 0.5 66.7 0.1 68.0 0.4 54.0 0.6
20.7 0.3 44.0 0.4 6.1 0.7 11.9 0.3 n.d. 18.9 0.1 22.6 0.1 38.2 0.1 17.5 0.1 29.5 0.1 16.5 1.3 59.0 1.1 10.9 0.8 52.1 0.5 6.9 0.8 34.0 0.9 52.7 0.4 18.9 0.7 59.6 0.6 63.1 0.3 81.9 0.4 68.5 0.2 91.6 0.1 70.3 0.1 94.0 0.2 4.3 0.5 6.4 0.5 n.d. 15.2 0.1 25.5 0.1 41.3 0.2 60.9 2.5 21.2 0.6 46.0 0.4 16.1 0.2 39.9 0.2 8.2 0.3 3.5 0.1 9.0 0.1 21.5 0.2 36.7 0.2 16.0 0.1 20.3 0.2 18.3 0.1 8.1 0.3 12.2 0.4
29.5 37.5 0.4 0.4 0.4 0.1 0.1 0.1
65.5 53.0 49.3 34.8 30.1 3.9 1.2 0.4
5.0 9.5 50.3 64.8 69.5 96.0 98.7 99.5
Diffusion couple
The primary phases and the solidification sequence were carefully analysed for all alloys. The microstructures of the as-cast alloys show different characteristics, as illustrated in Fig. 1(a)–(h). Of the alloys in the Ti-rich part (> 50 at.% Ti) the as-cast microstructures of alloys 01 and 03 show only the two phases b-Ti and (Fe,Ni)Ti, with b-Ti as the primary phase in alloy 01 and (Fe,Ni)Ti as the primary phase in alloy 03 (Fig. 1(a)). The as-cast microstructures of the alloys 04–07, 16, 32, 33 and 35 show a three-phase microstructure of b-Ti, (Ni,Fe)Ti2 and (Fe,Ni)Ti, with (Fe,Ni)Ti as the primary phase. However, they show two different phase arrangements, i.e. different solidification sequences. In alloys 04, 07 and 32 (Fig. 1(b)) primary dendrites of (Fe,Ni)Ti are embedded within b-Ti and only a few small precipitates of (Ni,Fe)Ti2 are observed. In contrast, alloys 05, 06, 16, 33 and 35 (Fig. 1(c)) show primary (Fe,Ni)Ti surrounded by (Ni,Fe)Ti2 which apparently formed via the peritectic reaction L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2. In between the (Ni,Fe)Ti2 grains some small regions with a lamellar eutectic microstructure composed of b-Ti and (Ni,Fe)Ti2 are found (encircled in black in Fig. 1(c)), which formed from the last liquid. Alloy 15 shows the phases b-Ti and (Ni,Fe)Ti2 in the as-cast state, with b-Ti as the primary phase (Fig. 1(d)).
‘‘Liquid’’ (Fe,Ni)Ti (Fe,Ni)Ti Fe2Ti Fe2Ti a-Fe a-Fe g-Fe n.d., not possible to determine.
Of the alloys in the Ti-lean part (<50 at.% Ti) alloys 11 (Fig. 1(g)), 20 (Fig. 1(e)), 21 show a eutectic microstructure or small regions of eutectic, as in the case of alloy 08 (Fig. 1(f)). These eutectic regions as well the phases involved were analysed by EPMA. The same eutectic composition was observed for alloys 20 and 21, although they show different primary phases. The primary phase in alloy 20 (Fig. 1(e)) is Fe2Ti (29.3Ni; 33.7Ti; 37.0Fe), which is surrounded by a eutectic microstructure of the composition 46.0Ni; 35.8Ti; 18.2Fe. The eutectic consists of the three phases (Fe,Ni)Ti (41.1Ni; 40.0Ti; 18.9Fe), Ni3Ti (51.6Ni; 33.1Ti; 15.3Fe) and Fe2Ti (37.8Ni; 34.4Ti; 27.8Fe).
L.I. Duarte et al. / Intermetallics 18 (2010) 374–384 Table 3 Compositions of coexisting phases measured by EPMA for alloys annealed at 800 C. Alloy no.
Phase
Ni (at.%)
Ti (at.%)
Fe (at.%)
01
b-Ti
0.7 0.1 5.1 0.1 0.6 0.1 4.7 0.1 2.1 0.1 6.3 0.1 4.1 1.1 9.8 0.1 8.0 0.3 15.0 0.1 16.5 0.5 17.7 60.2 30.7 68.2 37.1 72.7 57.4 12.9 58.5 32.1 8.9 0.1 11.1 0.1 3.3 0.1 2.2 0.1 1.0 0.1 0.6 0.1 5.2 0.1 24.9 0.1 17.3 0.2 22.4 0.1 8.1 0.3 3.7 0.3 30.9 17.5 50.6 31.4 18.6 52.0 38.9 0.4 69.2 3.1 10.0 0.5 2.1 0.1 11.6 0.1 15.0 0.3 25.1 0.3
82.3 0.3 67.3 0.2 82.6 0.1 67.2 0.2 51.7 0.3 67.1 0.1 52.0 0.4 67.2 0.1 52.2 0.9 67.4 0.1 52.4 0.4 28.5 22.6 10.7 22.2 5.8 20.8 3.6 27.6 21.6 11.5 29.3 0.2 4.6 0.5 28.7 0.7 3.9 0.7 29.0 0.4 3.2 0.2 91.7 0.1 68.4 0.1 67.5 0.3 52.1 0.3 49.1 0.8 35.9 1.0 47.8 34.5 34.1 46.5 34.1 32.0 49.1 0.3 28.3 2.2 66.7 0.2 86.0 0.2 68.0 0.3 66.7 0.3 66.6 0.3
17.0 0.3 27.6 0.2 16.8 0.1 28.1 0.2 46.2 0.3 26.6 0.1 43.9 0.7 23.0 0.1 39.8 0.6 17.6 0.1 31.1 0.2 53.8 17.2 58.6 9.6 57.1 6.5 39.0 59.5 19.9 56.4 61.8 0.2 84.3 0.5 68.0 0.7 93.9 0.7 70.0 0.4 96.2 0.2 3.1 0.1 6.7 0.1 15.2 0.2 25.5 0.2 42.8 1.1 60.4 1.3 21.3 48.0 15.3 22.1 47.3 16.0 12.0 0.6 2.5 0.9 23.3 0.1 11.9 0.2 20.4 0.4 18.3 0.3 8.3 0.3
03
04 05 06 08
09 10 11
12 13 14 15 16 18 19
20
21 33 34 35 36
(Ni,Fe)Ti2 b-Ti (Ni,Fe)Ti2 (Fe,Ni)Ti (Ni,Fe)Ti2 (Fe,Ni)Ti (Ni,Fe)Ti2 (Fe,Ni)Ti (Ni,Fe)Ti2 (Fe,Ni)Ti Fe2Ti Ni3Ti g-(Fe,Ni) Ni3Ti g-(Fe,Ni) Ni3Ti g-(Fe,Ni) Fe2Ti Ni3Ti g-(Fe,Ni) Fe2Ti g-(Fe,Ni) Fe2Ti a-Fe Fe2Ti a-Fe b-Ti (Ni,Fe)Ti2 (Ni,Fe)Ti2 (Fe,Ni)Ti (Fe,Ni)Ti Fe2Ti (Fe,Ni)Ti Fe2Ti Ni3Ti (Fe,Ni)Ti Fe2Ti Ni3Ti (Fe,Ni)Ti Fe2Ti (Ni,Fe)Ti2 b-Ti (Ni,Fe)Ti2 (Ni,Fe)Ti2 (Ni,Fe)Ti2
When no standard deviation is given for an analysis only few points could be measured indicating that the phase was either too small for a measurement or that only few particles could be detected.
The as-cast microstructures of the alloys 08 (Fig. 1(f)), 09 and 10 all show primary dendrites of g-(Fe,Ni) surrounded by grains of phases Ni3Ti and g-(Fe,Ni). In addition, in alloy 08 small eutectic regions were observed (marked by arrows in Fig. 1(f)). The as-cast microstructure of alloy 11 (Fig. 1(g)) also shows a eutectic microstructure. The composition of this eutectic is 31.5Ni; 20.9Ti; 47.6Fe and it consists of the three phases g-(Fe,Ni) (33.1Ni; 16.4Ti; 50.5Fe); Ni3Ti (35.0Ni; 21.8Ti; 43.2Fe) and Fe2Ti (25.4Ni; 24.3Ti; 50.3Fe). Alloys 13 and 14 both show microstructures where primary a-Fe is surrounded by a lamellar eutectic of a-Fe and Fe2Ti. In alloy 12 (Fig. 1(h)) primary Fe2Ti is surrounded by g-Fe and some more Fe2Ti, while in alloy 18 primary Fe2Ti is embedded in (Fe,Ni)Ti, which apparently formed via the peritectic reaction L þ Fe2Ti 4 (Fe,Ni)Ti. 3.2. Isothermal section at 1000 C After heat treatment at 1000 C the microstructures of all samples were characterized and the phase compositions were determined by EPMA. Equilibria between the phases b-Ti, (Ni,Fe)Ti2
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and (Fe,Ni)Ti as well as between (Ni,Fe)Ti2 and (Fe,Ni)Ti can be observed for the Ti-rich specimens. Fig. 2(a) shows the three-phase equilibrium in alloy 05, whereas the two-phase equilibrium observed in alloy 06 is presented in Fig. 2b. Two different three-phase equilibria were found for the two alloys presented in Fig. 3. Alloy 20 shows the equilibrium between (Fe,Ni)Ti, Fe2Ti and Ni3Ti (Fig. 3(a)). The other three-phase equilibrium between Fe2Ti, Ni3Ti and g-(Fe,Ni) is found in alloy 11 (Fig. 3(b)). In alloys 02 and 15 melt is present at 1000 C, though the annealing time of 10 min was too short to attain equilibrium. The diffusion couple was prepared in order to provide an overview of the phase reactions in the Ti-rich corner of the ternary system at 1000 C. Fig. 4 shows the schematic drawing of this diffusion couple as well as a backscattered electron image from its reaction zone. Strong interdiffusion and phase reactions occurred, resulting in the formation of a series of layers. The total thickness of these layers is about 25–30 mm. The phases (Ni,Fe)Ti2, (Fe,Ni)Ti, Fe2Ti and a-Fe were found in this reaction zone, as indicated in Fig. 4. Several EPMA line scans were performed across the diffusion zone, and from the resulting concentration profiles phase equilibria were determined by extrapolations to the phase boundaries. A tieline obtained from the ‘‘frozen zone’’ and (Ni,Fe)Ti2 is shown in Fig. 5. The other tie-lines follow the binary Fe-Ti phase diagram, being in good agreement with the binary phase compositions. The compositions of the coexisting phases as established by EPMA on samples quenched from 1000 C are shown in Fig. 5, and all compositions are compiled in Table 2. 3.3. Isothermal section at 800 C After heat treatment at 800 C a three-phase microstructure was found in four alloys on the Ti-rich side. Fig. 6(a) presents an example of the three-phase microstructure of alloy 03, showing the equilibrium between b-Ti, (Fe,Ni)Ti and (Ni,Fe)Ti2. For the same temperature an equilibrium between (Ni,Fe)Ti2 and (Fe,Ni)Ti was found in alloys with higher Ni content, for example for the alloys 05, 06, and 16. In Fig. 6(b) the two-phase microstructure of alloy 06 is presented. The equilibrium b-Ti þ (Ni,Fe)Ti2 was found for the alloys 01, 15 and 34. As an example for this equilibrium the microstructure of alloy 15 is shown in Fig. 7(a). Two different three-phase equilibria were found for the alloys with Ti contents <50 at.%. Alloys 19 and 20 show the equilibrium between (Fe,Ni)Ti, Fe2Ti and Ni3Ti. As an example, the microstructure of alloy 19 is presented in Fig. 7(b). The second three-phase equilibrium is between Fe2Ti, Ni3Ti and g-(Fe,Ni). This was observed in the alloys 08 and 11, Fig. 8(a). The alloys 09 and 10 show two-phase equilibria between the Ni3Ti and g-(Fe,Ni) phases. This equilibrium is represented by alloy 09 in Fig. 8b. As can be seen, these samples have a fine lamellar structure that is difficult to analyse by EPMA. The compositions of the coexisting phases as established by EPMA on samples quenched from 800 C are shown in Fig. 9, and all compositions of the phases are compiled in Table 3. 3.4. DTA results As mentioned above, only the data from the first heating curve were considered because of the high reactivity of Ti with the crucible. The yttria coating of the alumina crucible helped to prevent a strong reaction. However, complete suppression of any reaction between the crucible and the melt was not possible for all specimens. The first onset temperatures of melting, which were estimated for each curve from the intersection of the extrapolated baseline with the back-extrapolated peak, are presented in Table 1. These temperatures represent the solidus temperature in all cases.
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Fig. 1. As-cast microstructures of: (a) alloy 03: Fe31.4Ni2.0Ti66.6; (b) alloy 32: Fe28.4Ni5.0Ti66.6; (c) alloy 35: Fe18.3Ni15.0Ti66.7; (d) alloy 15: Fe4.6Ni14.8Ti80.6; (e) alloy 20: Fe22.5Ni43.5Ti34.0; (f) alloy 08: Fe39.0Ni44.4Ti16.6; (g) alloy 11: Fe50.0Ni29.6Ti20.4 and (h) alloy 12: Fe67.8Ni9.4Ti22.8.
Fig. 10 shows the DTA curves recorded during heating of homogenized alloys with a fixed Ti content of 66.7 at.%. The results of these samples are presented in order to obtain a better understanding of the formation of (Ni,Fe)Ti2. It is observed that the solidus temperature decreases monotonously with increasing Ni content to a value of about 1020 C 3 C for the alloys between 10.0 and 15.0 at.% Ni.
For 15.8 at.% Ni a slightly higher onset temperature of 1024 C 3 C is seen, while for the specimen with the highest Ni content of 25.1 at.% an onset temperature of 1018 C 3 C is observed. The onset temperature of 1028 C 3 C found for the alloys with a Ni content of 6.5 at.% suggests that an invariant reaction takes place at about this composition; see Section 4.2.
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Fig. 2. (a) Three-phase microstructure of alloy 05: Fe23.2Ni10.3Ti66.5 and (b) two-phase microstructure of alloy 06: Fe17.7Ni15.8Ti66.5 after annealing at 1000 C for 1167 and 375 h, respectively.
Fig. 3. Three-phase microstructure of two alloys after annealing at 1000 C for 1167 h. (a) Alloy 20: Fe22.5Ni43.5Ti34.0; (b) alloy 11: Fe50.0Ni29.6Ti20.4.
The liquidus temperature could be determined for some alloys by extrapolating the end of the last peak before the return to the baseline, as shown in Fig. 10 for two examples. Unfortunately, in some cases this effect is not well-defined, or the experiment was stopped before melting was completed in order to prevent reactions with the crucible. The determination of the liquidus temperatures is shown in an example in Fig. 10. Liquidus temperatures were only derived from experiments where no reactions with the crucible were observed on the DTA curves and where, by optical inspection after the experiment, no reaction was noticed between sample and crucible. The liquidus temperatures are also given in Table 1.
g-(Fe,Ni) was extensively discussed by Gupta [1], who compared the data with the work done by Drake [16] in which the maximum solubility of Ti in g-(Fe,Ni) was found to be as high as 12–13 at.%. Gupta [1] used the latter data for the boundary of the g-(Fe,Ni) phase, also taking into consideration the work done by Abramycheva et al. [11]. The results of the tie-lines obtained in this work and the resulting boundary for the g-(Fe,Ni) phase, as well as the position of the three-phase field Fe2Ti þ Ni3Ti þ g-(Fe,Ni), are in good agreement with the data assessed by Gupta [1]. It may be speculated that coherency stresses are responsible for the different
4. Discussion 4.1. Isothermal sections at 1000 and 800 C The tie-lines and tie-triangles obtained from evaluation of the individual alloys heat-treated at 1000 C and 800 C are shown in Figs. 5 and 9 and the isothermal sections established from these data are shown in Figs. 11 and 12, respectively. The observation that the compositions of the alloys fall onto the tie-lines or lie within the tie-triangles confirms that the actual compositions are very close to those intended. The results obtained for 1000 C from the alloys with lower Ti content show good agreement with the tie-lines from literature [11,14], as can be seen in Fig. 5. It is noted that some investigators [6,12,15] reported a significantly lower solubility limit of Ti in g-(Fe,Ni). This scattering of the solubility limit of Ti in the phase
Fig. 4. Diffusion couple (schematic illustration and BSE image of reaction zone) between Fe and NiTi2 annealed for 30 min at 1010 C in a vacuum furnace.
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Fig. 5. Isothermal section of the Fe–Ni–Ti system at 1000 C showing EPMA data measured on individual alloys and with experimental data from literature. Empty squares represent the phase compositions according to the accepted assessments of the binary systems.
data reported for the solid solubility of Ti in g-(Fe,Ni), as these stresses may vary considerably in dependence on annealing time and/or cooling rate. According to the present results, the solid solubility for Ti in g-(Fe,Ni) decreases with decreasing temperature (c.f. Figs. 11 and 12) in accordance with the latest assessment [4] and as expected from the data in [7,16]. The maximum solid solubility for Ti in g-(Fe,Ni) is found at the three-phase equilibrium g-(Fe,Ni) þ Fe2Ti þ Ni3Ti. This result also agrees with the experimentally determined isotherms for 900 C [7] and 1027 C [16]. However, the present data contradict earlier data from Speich [6], according to which the solid solubility for Ti in g-(Fe,Ni) increases continuously with increasing Ni content at 1000 C, and data by Abramycheva et al. [11], according to which the solid solubility for Ti in g-(Fe,Ni) at 1000 C has its maximum not at the three-phase equilibrium but at a Ni content which is about 10 at.% higher. Therefore, qualitatively the data confirm that the solubility of Ti in g-(Fe,Ni) varies in dependence on the Fe:Ni ratio according to what was found at 1027 C [16] and at 900 C [7].
Quantitatively, the maximum solid solubility of Ti in g-(Ni,Fe) is given by the results for alloys 11 and 08 at 800 C, which yield 11.5 and 10.7 at.% Ti, respectively (see Table 3), and the maximum value at 1000 C of 12.0 at.% Ti is found for alloy 11. The main differences between this work and the isothermal section at 1000 C proposed by Abramycheva et al. [11] exist in the Ti-rich part, i.e. at a Ti content > 50 at.%. In this part Abramycheva et al. [11] investigated only one diffusion couple between the binary alloys Fe21.7Ni78.3 and Fe43.5Ti56.5 (in at.%). In this diffusion couple they found the phase equilibria (Fe,Ni)Ti þ b-Ti and (Fe,Ni)Ti þ (Ni,Fe)Ti2 from which they deduced the position of the tietriangle (Fe,Ni)Ti þ b-Ti þ (Ni,Fe)Ti2. All other phase equilibria given by Abramycheva et al. [11] in the Ti-rich part are given by dashed lines, indicating that they are tentative until experimental evidence exists. The three-phase equilibrium (Fe,Ni)Ti þ b-Ti þ (Ni,Fe)Ti2 was also found in the present investigation, though a lower Fe content for (Ni,Fe)Ti2 was measured. However the microstructures of several alloys are not in agreement with the isothermal section
Fig. 6. (a) Three-phase microstructure of alloy 03: Fe31.4Ni2.0Ti66.6 and (b) two-phase microstructure alloy 06: Fe17.7Ni15.8Ti66.5 after annealing at 800 C for 1119 and 375 h, respectively.
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Fig. 7. (a) Two-phase microstructure of alloy 15: Fe4.6Ni14.8Ti80.6 and (b) three-phase microstructure of alloy 19: Fe32.8Ni27.2Ti40.0 after annealing at 800 C for 375 and 1100 h, respectively.
Fig. 8. (a) Three-phase microstructure of alloy 08: Fe39.0Ni44.4Ti16.6 and (b) two-phase microstructure alloy 09: Fe29.6Ni55.4Ti15.0 after annealing at 800 C for 1119 h.
proposed in Ref. [11]. For example, according to [11] the alloys 05, 06 and 16 should show a two-phase equilibrium between b-Ti and (Fe,Ni)Ti. Instead, an equilibrium between (Ni,Fe)Ti2 and (Fe,Ni)Ti was found, as can be seen in Fig. 2(b). This means that the phase (Ni,Fe)Ti2 has a much wider homogeneity range at 1000 C than assumed in [11]. In consequence, a second three-phase equilibrium (Fe,Ni)Ti þ bTi þ (Ni,Fe)Ti2 and the three-phase equilibrium (Fe,Ni)Ti þ b-Ti þ L, which were both proposed by Abramycheva et al. [11], do not exist. Instead, the two three-phase equilibria L þ b-Ti þ (Ni,Fe)Ti2 and L þ (Fe,Ni)Ti þ (Ni,Fe)Ti2 exist at 1000 C. Though these three-phase equilibria were not observed in any of the present samples because they include the liquid phase, the position of the respective tietriangles can be localised by the adjoining two-phase equilibria that were experimentally determined. For L þ (Fe,Ni)Ti þ (Ni,Fe)Ti2 the composition of L should be at a slightly higher Fe content than that determined from the diffusion couple (c.f. Fig. 4), and the composition of (Ni,Fe)Ti2 should be at a slightly lower Fe content than that observed for this phase in alloy 36, while the composition of (Fe,Ni)Ti must lie between the compositions of (Fe,Ni)Ti determined in alloy 36 and in the diffusion couple. The phase compositions of the three-phase equilibrium L þ b-Ti þ (Ni,Fe)Ti2 are more difficult to fix. The composition of b-Ti should include the two compositions determined for this phase in alloys 02 and 15, and the composition of (Ni,Fe)Ti2 should be at a slightly higher Fe content than that determined in alloy 02. For the composition of L no data exist, and it can therefore only be estimated. Because both threephase equilibria were not observed directly they are given as dashed lines in Fig. 11.
So far no isothermal section at 800 C is available in literature. The isothermal section in Fig. 12 was based on the current EPMA results shown in Fig. 9. Even after prolonged annealing at this
Fig. 9. Compositions of alloys (solid black circles) and compositions of phases as determined by EPMA (empty black circles) of the Fe–Ni–Ti system at 800 C.
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Fig. 10. DTA curves of the Ti-rich alloys with 66.7 at.% of Ti obtained during heating at 10 K/min. Onset temperatures of melting (left) and liquidus temperatures (right) are indicated.
temperature, it is difficult to generate phases that are large enough for EPMA measurements, as can be seen in Figs. 3(b) and 8 for the Ni3Ti phase. Therefore, concentrations in the Ti-lean part of the system are somewhat uncertain and are indicated by dashed lines. The fine lamellar structure observed in Fig. 8 for both alloys is due to the ‘‘cellular precipitation reaction’’ described by Speich [6] and Fournelle [12]. According to these reports the decomposition of Fe-30Ni-6Ti (in wt.%) occurs via a cellular precipitation reaction in which cells consisting of alternate lamellae of Ni3Ti (DO24) and austenite are formed. This first cellular reaction is subsequently consumed by a second cellular reaction of much coarser spacing. These authors also reported that Ni3Ti forms an almost perfectly coherent interface with the austenite, with the following orientation relationship:
gB þ Ni3Ti / gC þ (Fe,Ni)2Ti, where gA, gB and gC are austenitic solid solutions of different solute content [6]. As well as in Fig. 8 these features can also be observed in Fig. 3(b), where lamellar Ni3Ti grew inside of the g-(Fe,Ni) grains. No ternary compound could be found at both temperatures (Figs. 11 and 12). Compared to the section at 900 C by van Loo et al. [7], the same phase equilibria were observed at 800 C, with all binary compounds showing a large solubility for the third element. From 900 C to 1000 C, phase equilibria in the Ti-rich corner change because above 942 C a melt forms within the Ni–Ti binary system.
ð001ÞNi3 Ti ==ð111Þg; ½010Ni3 Ti ==½110g
From the examination of the as-cast microstructures and the DTA investigations the liquidus projection shown in Fig. 13 was established. In the Ti-rich part, (Fe,Ni)Ti was observed as a primary phase in all alloys with 66.7 at.% Ti which contained between 2 and 25.1 at.% Ni (alloys 03–06, 32, 33, 35, 36). For the alloys with up to 5.0 at.% Ni (alloys 03, 04, 32), primary dendrites of (Fe,Ni)Ti are surrounded by b-Ti (Fig. 1(a and b)) although only in the case of alloy 03 is there no indication of any reaction between these two phases. In addition, alloys 04 and 32, i.e. those with the highest Ni content, show small amounts of eutectic (Ni,Fe)Ti2 þ b-Ti (Fig. 1(b)) which should be formed from the last solidifying melt. Apparently the (Ni,Fe)Ti2 that surrounds the (Fe,Ni)Ti, (Fig. 1(b)) is already formed in alloys 04 and 32 by the peritectic reaction L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2. The latter feature was also observed by [18]. In contrast, in the alloys with Ni content of between 10.0 and 15.8 at.% Ni (05, 06, 35, 33) (Fe,Ni)Ti is surrounded by (Ni,Fe)Ti2, with clear evidence that (Fe,Ni)Ti have partly dissolved in order to form (Ni,Fe)Ti2 (Fig. 1(c)), i.e. (Ni,Fe)Ti2 formed via the peritectic reaction L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2. The above observations and also those made for the other alloys in the Ti-rich part are in full agreement with the liquidus projection proposed by Riani et al. [18], and the Ti-rich part in Fig. 13 was drawn accordingly. In addition, the temperatures for the invariant reaction U2: L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 þ b-Ti (at about 6 at.% Ni) of about 1030 C and the maximum (M2) on the monovariant line L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 at approximately 1035 C established here agree well with the data assessed by Cacciamani et al. [3], who gave w1025 C and w1030 C, respectively. It is noted that a recent thermodynamic modelling of this system published by
It was reported by Speich [6] that the ternary alloys may exhibit two types of discontinuous precipitation: 1) gA / gB þ Ni3Ti and 2)
Fig. 11. Isothermal section of the Fe–Ni–Ti system at 1000 C. Empty squares represent the phase compositions according to the accepted assessments of the binary systems.
4.2. Liquidus surface and reaction scheme
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Fig. 12. Isothermal section of the Fe–Ni–Ti system at 800 C. Empty squares represent the phase compositions according to the accepted assessments of the binary systems.
De Keyzer et al. [19] yields slightly higher temperatures for these invariant reactions (U2 ¼ 1043 C and M2 ¼ 1044 C). In contrast to the Ti-rich part the solidification behaviour of the Ti-lean alloys was found to differ from liquidus projections established previously, and therefore this part will be discussed in more detail. The solidification of alloys 13 and 14 both end at the eutectic trough e2–U1, and the Ni-rich alloy (13) at a somewhat lower temperature, which is in full agreement with the run of the eutectic trough, which descends from 1289 C (e2) in the Fe–Ti binary system to about 1199 C at the invariant reaction U1. The temperature for U1 was not observed directly but is estimated from the onset temperature observed for alloy 12, which is 1199 C (Table 1). From the composition of alloy 12 and the composition of the primary Fe2Ti phase (7.9Ni; 27.8Ti; 64.3Fe) in that alloy we conclude that the effect at 1199 C belongs to the reaction U1. This temperature is in good agreement with that of 1200 C given in the last assessments [3,4], although the calculated value (1270 C) [19] is much higher. While onset temperatures and metallographic observations for alloys 13 and 18 confirm that the peritectic reaction L þ Fe2Ti 4 (Fe,Ni)Ti continues from p2 with decreasing temperature into the ternary system, observations for alloys 20 and 21 contradict previous liquidus projections [1,3,4,18]. According to these previous works the solidification of alloys 20 and 21 will end up at the eutectic trough e3–U2 (labelled e3–E2 in Fig. 13; see below), whose temperature was supposed to descend towards to the binary system from U2 1118 < T < 1317 C to e3 1118 C. Since the first onsets for both alloys are below 1118 C and because the Ni-rich alloy (21) has the higher onset (1113 C), this means that the eutectic trough descends into the ternary and not towards the binary. Moreover, the onset for alloy 20 is at 1099 C, which is the lowest temperature observed for all the Ti-lean alloys. That the eutectic trough descends inward and that the lowest onset temperature is observed here consequently means that the invariant reaction U2 is actually a ternary eutectic (E2: L 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti). This conclusion is also confirmed by the EPMA results which show that a eutectic is located at 46.0Ni; 35.8Ti; 18.2Fe and that it consists of the phases Fe2Ti, Ni3Ti and (Fe,Ni)Ti.
383
Fig. 13. Proposed liquidus projection of the Fe–Ni–Ti system according to this study.
In view of the large volume fraction of this ternary eutectic (Fig. 1(e)) the onset temperature of 1099 C for alloy 20 should actually be the temperature of E2. This temperature is in good agreement with that recently calculated by De Keyzer et al. (E2 ¼ 1103 C) [19], although the calculated composition is at a much higher Ni content than that established by EPMA. The DTA results for alloys 08–10 (Table 1) show that the onset temperatures decrease from alloy 10 to 08, agreeing with previous work. The solidification of alloys 09 and 10 ends up at the eutectic trough e1–E1, with the Ni-rich alloy (10) showing a somewhat higher solidification temperature (1267 C). This agrees fully with the run of the eutectic trough, which descends from 1304 C from the binary eutectic e1 towards E1. According to previous assessments [1,3,4,18] and the calculated liquidus projection [19] alloy 08 should end up at the ternary eutectic E1 at a temperature of around 1110 C and be the closest alloy to this reaction. In fact, the microstructure of the as-cast alloy 08 shows only small amounts of eutectic regions (Fig. 1(f)). That the signal observed in the DTA for E1 in alloy 08 (at 1109 C) is only small is consistent with this observation. Actually, EPMA investigations on alloy 11 revealed that the composition of the eutectic E1 (31.5Ni; 20.9Ti; 47.6Fe) is close to this alloy, i.e. at a considerable lower Ni content than previously assumed. According to the first onset observed for alloy 11 in the DTA the temperature for E1 is 1108 C. Between the two ternary eutectics there must be a maximum in temperature on the invariant line L þ Fe2Ti þ Ni3Ti. According to Alkemade’s theorem [20] this maximum should be at the point where the connecting line between the congruent compositions of Fe2Ti and Ni3Ti (so-called Alkemade line) intersects with the invariant line. This point is represented by M1 in Fig. 13 and its temperature is estimated from the course of the isotherms to be at about 1150 C, which is a little bit lower than the temperature of 1187 C calculated by De Keyzer et al. [19]. In view of the above, the liquidus projection as well as the reaction scheme must be modified as follows: a) U2 becomes E2: L 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti at about 1099 C; b) E1: L 4 g(Fe,Ni) þ Fe2Ti þ Ni3Ti at 1108 C; c) the line E1–E2 has a maximum (M1) at around 1150 C. It is noted that the modified liquidus
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Fig. 14. Reaction scheme of the Fe–Ni–Ti system.
projection (Fig. 13) and the reaction scheme (Fig. 14) are consistent with one another. 5. Conclusions Isothermal sections of the ternary system Fe–Ni–Ti were established at 1000 C and for the first time at 800 C. A good correlation was found between our results and those from literature for both isothermal sections. Qualitatively the present data confirm that the solid solubility of Ti in g-(Fe,Ni) a) varies in dependence on the Fe:Ni ratio, as found before at 1027 C [16] and at 900 C [7]; and b) decreases with decreasing temperature (c.f. Figs. 11 and 12). The liquidus projection as well as the reaction scheme in the Tilean part must be modified as follows: a) U2 becomes E2: at 1099 C; b) E1: L 4 gL 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti (Fe,Ni) þ Fe2Ti þ Ni3Ti at 1108 C; c) the line E1–E2 has a maximum (M2) at around 1150 C. The invariant reaction in the Ti-rich part appears to be of quasiperitectic four-phase transition-type (U) L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 þ bTi at around 1030 C. This implies the existence of a maximum (M1) on the monovariant liquidus line, at approximately 1035 C, as already suggested by Cacciamani et al. [3] and Riani et al. [18]. Further, a three-phase equilibrium between b-Ti, (Fe,Ni)Ti and (Ni,Fe)Ti2 was observed in some alloys, but neither the three-phase equilibrium b-Ti þ L þ (Fe,Ni)Ti, nor the two-phase equilibrium bTi þ (Fe,Ni)Ti were found in any sample. Instead, an equilibrium between (Ni,Fe)Ti2 and (Fe,Ni)Ti was observed. Alternative liquidus projection and reaction schemes were proposed based on these new experimental data.
Acknowledgements This work was performed within the framework of the COST Action 535, ‘‘Thermodynamics of alloyed aluminides (THALU)’’, and was financially supported by the Switzerland State Secretariat for Education and Research (SER) under Contract no. C05.0031.
References [1] Gupta KP. Phase diagrams of ternary nickel alloys, part 1. Indian Inst Metals 1990:321–43. [2] Gupta KP. J Phase Equilib 2001;22(2):171–5. [3] Cacciamani G, De Keyzer J, Ferro R, Klotz UE, Lacaze J, Wollants P. Intermetallics 2006;14:1312–25. [4] Ghosh G. FeNiTi (ironnickeltitanium), part 4. In: Effenberg G, Ilyenko S, editors. Light metal systems. Heidelberg: Springer Berlin; 2006. p. 299–316. [5] Vogel R, Wallbaum HS. Arch Eisenhuttenwesen 1938;12:299–304. [6] Speich GR. Trans Met Soc AIME 1963;227:754–62. [7] van Loo FJJ, Vrolijk JWGA, Bastin GF. J Less Common Met 1981;77:121–30. [8] Dudkina LP, Kornilov II . Russ Metall 1967;4:98–101. [9] Alisova SP, Budberg PB, Barmina TI, Lutskaya NV. Russ Metall 1994;1:117–21. [10] Efimenko LP, Petrova LP, Sviridov SI. Russ Metall 2000;3:42–6. [11] Abramycheva NL, V’yunitskii IV, Kalmykov KB, Dunaev SF. Vestn Mosk Univ Ser 2 Khim 1999;40:139–43. [12] Fournelle RA. Acta Metall 1979;27:1135–45. [13] Abraham JK, Pascover JS. Trans Met Soc AIME 1969;245:759–68. [14] Jia CC, Ishida K, Nishizawa T. Metall Mater Trans A 1994;A25:473–85. [15] Abraham JK. Trans Met Soc AIME 1968;242:2365–7. [16] Drake S.C. B.Sc. thesis, MIT; 1982. p. 34. [17] Sugiura T, Ishikawa S, Matsuo T, Takeyama M. Mater Sci Forum 2007;561– 565:435–8. [18] Riani P, Cacciamani G, Thebault Y, Lacaze J. Intermetallics 2006;14:1226–30. [19] De Keyzer J, Cacciamani G, Dupin N, Wollants P. Calphad 2009;33:109–23. [20] Ehlers EG. The interpretation of geological phase diagrams. San Francisco: Freeman, W.H. and Company; 1972.
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Experimental study of the Fe–Ni–Ti system L.I. Duarte a, b, *, U.E. Klotz a,1, C. Leinenbach a, M. Palm c, F. Stein c, J.F. Lo¨ffler b a
¨ berlandstrasse 129, CH-8600 Du ¨ bendorf, Switzerland Empa, Swiss Federal Laboratories for Materials Testing and Research, Laboratory for Joining and Interface Technology, U Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, CH-8093 Zurich, Switzerland c ¨ r Eisenforschung GmbH, Max-Planck-Str. 1, D-40237 Du ¨ sseldorf, Germany Max-Planck-Institut fu b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 8 August 2008 Received in revised form 12 August 2009 Accepted 13 August 2009 Available online 3 September 2009
In this investigation phase relations in the Fe–Ni–Ti system were studied and two isothermal sections at 800 C and 1000 C as well as a revised liquidus projection were established. Microstructural characterisation of the as-cast alloys and of samples equilibrated at 800 and 1000 C was performed by scanning electron microscopy (SEM), chemical compositions of the phases were analysed by electron probe microanalysis (EPMA), and liquidus temperatures were examined by differential thermal analysis (DTA). The experimental results clarify some uncertainties concerning the melting behaviour and the solid-state phase equilibria between the phases (Ni,Fe)Ti2, (Fe,Ni)Ti and b-Ti. The present data also confirm that the solid solubility of Ti in g-(Fe,Ni) varies in dependence on the Fe:Ni ratio and decreases with decreasing temperature. The liquidus projection as well as the reaction scheme in the Ti-lean part are modified because two ternaries eutectic E1: L 4 g-(Fe,Ni) þ Fe2Ti þ Ni3Ti and E2: L 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti are found at 1108 and 1099 C, respectively. Ó 2009 Elsevier Ltd. All rights reserved.
Keywords: A. Ternary alloy systems A. Laves phases B. Phase diagrams F. Microprobe
1. Introduction Fe–Ni–Ti is one of the ternary subsystems of the Al–Fe–Ni–Ti quaternary system, which is relevant for many different materials such as steels, Ni-based superalloys, alloyed aluminides, Ti-based corrosion-resistant alloys, shape-memory alloys, and amorphous materials. The Fe–Ni–Ti system has been studied extensively by a large number of investigators and was reviewed by Gupta [1] and later revised by the same author in Ref. [2]. More recently, this system was reviewed by Cacciamani et al. [3] and by Ghosh [4]. The liquidus projection and two isothermal sections at 900 C and 1000 C are known from literature, as discussed below. Gupta [1] proposed a schematic liquidus projection based mainly on the work of Vogel and Wallbaum [5], Speich [6], and Van Loo et al. [7]. The first comprehensive study of phase equilibria from the Fe–Ni–Ni3Ti– Fe2Ti part of this system was carried out in Ref. [5]. Thermal analysis and microstructural investigations were performed on several
* Corresponding author. Empa, Swiss Federal Laboratories for Materials Testing ¨ berlandstrasse 129, and Research, Laboratory for Joining and Interface Technology, U CH-8600 Du¨bendorf, Switzerland. Tel.: þ41 44 823 42 28; fax: þ41 44 823 40 39. E-mail address: [email protected] (L.I. Duarte). 1 Present address: FEM, Research Institute for Precious Metals and Metals Chemistry, Department of Physical Metallurgy, Katharinenstrasse 17, D-73525 Schwa¨bisch Gmu¨nd, Germany. 0966-9795/$ – see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2009.08.008
alloys and the results were reported in vertical sections (isopleths). From these results it was possible to draw a partial liquidus projection and several vertical sections. Later, Dudkina and Kornilov [8] established a vertical section TiFe–TiNi using twenty-one ternary alloys. They determined the solidus along this section and found complete solubility of Fe and Ni in the ordered phase (Fe,Ni)Ti. Alisova et al. [9] provided three vertical sections in the Tirich corner (>50 at.% Ti) for the Fe:Ni atomic ratios 3:1, 1:1, and 1:3 and a partial reaction scheme of the invariant equilibria in the Ti– TiNi–TiFe region. According to this work a four-phase transitiontype (U) reaction L þ (Fe,Ni)Ti 4 b-Ti þ NiTi2 takes place at 960 C and very close to the binary eutectic L 4 b-Ti þ NiTi2. The vertical sections provided in [9] are in disagreement with the results reported by Efimenko et al. [10]. In that work, diffusion couples were used which were produced by melting a mixture of Ni and Ti powders on an iron substrate and annealing at 1100, 1200 and 1250 C for 5–120 min. In the solidified melts, the two phases NiTi2 and (Fe,Ni)Ti were always observed, although according to Alisova et al. [9] an equilibrium between b-Ti and (Fe,Ni)Ti would have been expected. In the recent assessment by Gupta [2] the three vertical sections provided by Alisova et al. [9] were taken into consideration. However, the reaction L þ (Fe,Ni)Ti 4 b-Ti þ NiTi2 was shifted towards a much lower Fe content than in a previous assessment. Nevertheless, it was not stressed that the results from [9] disagree with the isothermal section at 900 C provided by van Loo et al. [7], who showed NiTi2 to have a high solubility for Fe at
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900 C. Ref. [7] provides the complete isothermal section at 900 C, showing that no ternary compound exists at this temperature and that all the binary compounds, i.e. NiTi2, NiTi, FeTi, Fe2Ti and Ni3Ti, possess a significant solubility for the third element. Abramycheva et al. [11] determined an isothermal section at 1000 C based on experimental results from several diffusion couples. In addition, a number of data have been reported for a Ti content of less than 50 at.% [6,12–14]. Speich [6] and Fournelle [12] studied the cellular precipitation of Ni3Ti in g-(Fe,Ni) in alloys with 28.7 at.% Ni and 7.0 at.% Ti. Speich [6] found a solvus at 982 C and a lower solubility limit of Ti in the g-(Fe,Ni) phase compared to Ref. [11]. Abraham [13,15] also found a lower solubility limit of Ti in the g-(Fe,Ni) phase determined at 1024 C in an alloy containing 27 at.% Ni. The scattering of the values for the solubility of Ti in the g-(Fe,Ni) phase was previously discussed in Ref. [1] and compared with the work done by Drake [16], where it was found that the maximum solubility of Ti in g-(Fe,Ni) at 1027 C was higher than 12–13 at.%. Consequently, the boundary of the g-(Fe,Ni) phase field was redrawn taking into consideration the work of Ref. [11]. It is noted that phase equilibria between g-(Fe,Ni), a-Fe and Fe2Ti (and L) at 1100 and 1200 C were recently determined by Sugiura et al. [17]. The isothermal section at 1000 C appears to be similar to that at 900 C for Ti content lower than 50 at.%. However, a number of differences exist in the Ti-rich part: from the binary Ni–Ti system the liquid phase extends into the ternary system at 1000 C. Also, according to [10], the (Ni,Fe)Ti2 phase, found in connection with very high Fe content, shows a very small homogeneity range at 1000 C. In addition, the nature and the position of the invariant reaction between the melt, (Ni,Fe)Ti2 and (Fe,Ni)Ti remains unclear. According to [11] the stability of (Ni,Fe)Ti2 at 1000 C is limited to an area close to Fe–Ti. An equilibrium between the liquid phase and (Ni,Fe)Ti2 was not observed. Instead, a three-phase equilibrium between b-Ti, (Ni,Fe)Ti2 and (Fe,Ni)Ti appears to exist at this temperature. In consequence, the field of primary solidification of (Ni,Fe)Ti2 must be substantially extended and the invariant reaction is not close to the Ni–Ti side but to the Fe–Ti side. If this is the case, two reactions are possible: an invariant peritectic (P) reaction with L þ b-Ti þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 or a ‘‘quasi-peritectic’’ four-phase transition-type (U) reaction with L þ (Fe,Ni)Ti 4 b-Ti þ (Ni,Fe)Ti2 associated with a maximum in the monovariant liquidus line connecting it with the binary peritectic L þ NiTi 4 NiTi2, as suggested in Ref. [3]. Riani et al. [18] chose the second option to explain their recent experimental results, obtained from DTA measurements and isothermal annealing at 900 C. The objective of this study is to determine the isothermal sections at 800 C and 1000 C experimentally and to clarify the uncertainties in the Ti-rich corner of this system. The determination of phase equilibria and solidification sequence at 66.7 at.% Ti is also an objective, to find out how far NiTi2 is stabilized by Fe. The solidus and liquidus temperatures of the Ti-rich alloys will provide a basis for a more complete thermodynamic description of the Fe–Ni–Ti system. 2. Experimental Pure elements of Fe (99.97%), Ni (99.98%) and Ti (99.99%) were used to produce the alloys given in Table 1. Alloys of about 4 g, most of them in the Ti-rich corner, were prepared by arc-melting under a purified argon atmosphere (99.999%) using a non-consumable tungsten electrode. All alloys were melted five times and inverted after three meltings to ensure homogeneity. In order to have a clean atmosphere during the melting, a Ti alloy was melted first as an oxygen getter. In addition, an oxygen cartridge in the argon line was used. No chemical analysis of the alloys was conducted since the weight losses during melting were less than 0.2 mass% in all cases. Afterwards, the samples were cut into 4 pieces for various
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Table 1 Nominal compositions of the Fe–Ni–Ti alloys produced by arc-melting, temperatures and annealing times of the heat treatments, and onset temperatures of melting and liquidus temperatures measured during heating of as-cast alloys by DTA. Nominal composition, (at.%) Alloy no.
Ni
01 02 03
2.7 74.1 23.2 800 19.5 69.4 11.1 1000 2.0 66.6 31.4 1000 800 6.5 66.6 26.9 800 10.3 66.5 23.2 1000 800 15.8 66.5 17.7 1000 800 4.6 58.4 37.0 – 44.4 16.6 39.0 1000 800 55.4 15.0 29.6 1000 800 65.3 14.0 20.7 1000 800 29.6 20.4 50.0 1000 800 9.4 22.8 67.8 1000 800 2.8 12.0 85.2 1000 800 0.9 13.9 85.2 1000 800 14.8 80.6 4.6 1000 800 19.2 61.7 19.1 1000 800 6.7 45.0 48.3 1000 800 27.2 40.0 32.8 1000 800 43.5 34.0 22.5 1000 800 56.9 36.9 6.2 1000 800 5.0 66.6 28.4 – 10.0 66.7 23.3 1000 9.0 72.2 18.8 1000 15.0 66.7 18.3 1000 25.1 66.6 8.3 1000
04 05 06 07 08 09 10 11 12 13 14 15 16 18 19 20 21 32 33 34 35 36
Ti
Fe
Temperature ( C)
DTA temperatures from heating ( C) Annealing time Onset Liquidus (h) temperature 1119 10 min 1167 1119 1119 1167 1119 375 375 – 1167 1119 1167 1119 1167 1119 1167 1119 1167 1119 1167 1119 1167 1119 10 min 375 375 375 1167 1100 1167 1100 1167 1100 1167 1100 – 702 702 702 702
1052 1128 988 1083 1062 1128 1028 1153 1020 1143 1024 1136 1044 1109 1252 1211 1267 1108 1179 1199 1270 1267 – 1282 – 975
–
1029 1160 –
–
1204 – 1099 1150 1113 1201 1040 1023 1022 1017 1018
1152 1158 1098 1159 1112
experiments. One piece was used to investigate the as-cast microstructure; another was used for differential thermal analysis (DTA) experiments; and the remaining two were used to investigate the microstructure after annealing experiments at 800 and 1000 C. For the heat treatments the samples were placed in evacuated quartz tubes re-filled with argon. Before the encapsulation the samples were cleaned in acetone, dried and wrapped in tantalum foil to avoid reactions between samples and the quartz tube. A titanium getter was placed in the quartz tube to absorb oxygen. The annealing experiments were performed at 800 C and 1000 C for 375 to 1200 h (c.f. Table 1) in an electric resistance furnace with a temperature accuracy of 3 C. After annealing, the samples were quenched in salt water to retain the equilibrium microstructures. In order to study the equilibrium between the liquid phase and NiTi2 at 1000 C a diffusion couple was prepared from an arcmelted alloy with the chemical composition of NiTi2 and a crucible of pure Fe (99.8%). The Fe crucible had a cylindrical shape of 20 mm diameter and 20 mm in height, with a hole of 16 mm in diameter and 17 mm in length. A NiTi2 sample of 16 mm in diameter and 5 mm in height was inserted into the Fe crucible. The diffusion couple was annealed for 30 min at 1010 C in a vacuum furnace (Torvac 1218H-PQ) under a pressure of <0.001 Pa.
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The heating rate was 10 C/min and the cooling rates were 50 C/min and 15 C/min from 1010 C to 800 C and from 800 C to room temperature, respectively. Higher cooling rates could not be achieved due to furnace limitations. After the annealing experiments, the samples were cut transversally by electro-discharge machining (EDM) and prepared by standard metallographic methods. The as-cast and equilibrated alloys were then examined by optical microscopy and scanning electron microscopy (SEM). The determination of the compositions of the coexisting phases of each sample was performed by electron probe microanalysis (EPMA) using a JEOL JXA8800 microanalyser with an acceleration voltage of 20 kV and a probe current of 2 108 A. For the measurement of the eutectic areas the beam was widened to 3–5 mm. Pure element standards provided by JEOL were used for calibration. The measurements have a relative accuracy of about 1%. If standard deviations given in Tables 2 and 3 are higher than this value or are missing only a few points could be measured, indicating that the phase was either too small to measure or that only a few particles could be detected. DTA experiments were performed in a Setaram SETSYS-18 and in a Netzsch STA 409 CD. Alumina crucibles coated with yttrium oxide were used in order to avoid reactions between the alloy and the crucible. The DTA cells were calibrated using the melting temperatures of the pure elements Al, Ag and Au. Cylindrical samples of 3 mm in diameter and a height of 5 mm, which were cut by EDM from equilibrated samples, were used in the DTA experiments. Before each experiment, the DTA cell was evacuated three times and re-filled with high-purity argon. Measurements were performed at a scanning rate of 10 C/min for heating and cooling under a continuous flow of argon. The evaluation of the DTA curves showed that the samples still exhibited some reactions with the crucibles due to the high reactivity of Ti in the liquid. Therefore only the data from the first heating curve were taken into account. 3. Results The overview in Table 1 shows the compositions of the alloys prepared by arc-melting, the temperatures and durations of the heat treatments, the onset temperatures of melting and the liquidus temperatures obtained from the DTA experiments. 3.1. As-cast alloys
Table 2 Compositions of coexisting phases measured by EPMA for alloys annealed at 1000 C and for the diffusion couple. Alloy no.
Phase
Individual alloys b-Ti 03 (Fe,Ni)Ti b-Ti 02 (Ni,Fe)Ti2 Liquid b-Ti 05 (Ni,Fe)Ti2 (Fe,Ni)Ti 06 (Ni,Fe)Ti2 (Fe,Ni)Ti 08 Ni3Ti g-(Fe,Ni) 09 Ni3Ti g-(Fe,Ni) 10 Ni3Ti g-(Fe,Ni) 11 Fe2Ti Ni3Ti g-(Fe,Ni) 12 Fe2Ti g-(Fe,Ni) 13 Fe2Ti a-Fe 14 Fe2Ti a-Fe b-Ti 15 (Ni, Fe)Ti2 Liquid 16 (Ni,Fe)Ti2 (Fe,Ni)Ti 18 (Fe,Ni)Ti Fe2Ti 19 (Fe,Ni)Ti Fe2Ti 20 (Fe,Ni)Ti Fe2Ti Ni3Ti 21 Ni3Ti (Fe,Ni)Ti 33 (Ni,Fe)Ti2 (Fe,Ni)Ti b-Ti 34 (Ni,Fe)Ti2 35 (Ni,Fe)Ti2 36 (Ni,Fe)Ti2 (Fe,Ni)Ti
Ni (at.%)
Ti (at.%)
Fe (at.%)
1.2 0.1 2.9 0.1 8.6 0.7 19.9 0.3 n.d. 3.8 0.1 10.2 0.1 9.4 0.1 15.1 0.1 18.1 0.1 59.7 0.9 30.7 0.7 66.0 0.5 40.8 0.4 71.1 0.4 59.7 0.3 18.9 0.3 56.8 0.7 28.4 0.4 8.7 0.1 11.2 0.1 3.0 0.1 2.7 0.1 0.9 0.1 0.7 0.1 8.8 0.5 24.5 0.5 n.d. 17.3 0.1 22.4 0.1 8.6 0.1 3.3 0.7 32.4 0.3 19.2 0.2 40.9 0.1 26.0 0.1 65.4 0.3 70.2 0.1 46.2 0.1 10.8 0.1 10.2 0.1 4.1 0.1 11.8 0.4 15.0 0.1 23.9 0.3 33.8 0.3
78.1 0.4 53.1 0.4 85.3 0.4 68.2 0.4 n.d. 77.3 0.1 67.2 0.3 52.4 0.1 67.4 0.3 52.4 0.1 23.8 0.5 10.3 0.4 23.1 0.4 7.1 0.2 22.0 0.4 6.3 0.6 28.4 0.6 24.3 0.1 12.0 0.2 28.2 0.3 6.9 0.4 28.5 0.2 5.7 0.1 28.8 0.1 5.3 0.2 86.9 0.3 69.0 0.4 n.d. 67.5 0.1 52.1 0.1 50.1 0.2 35.8 1.8 46.4 0.3 34.8 0.2 43.0 0.2 34.1 0.1 26.4 0.2 26.3 0.1 44.8 0.1 67.7 0.2 53.1 0.1 79.9 0.2 67.9 0.5 66.7 0.1 68.0 0.4 54.0 0.6
20.7 0.3 44.0 0.4 6.1 0.7 11.9 0.3 n.d. 18.9 0.1 22.6 0.1 38.2 0.1 17.5 0.1 29.5 0.1 16.5 1.3 59.0 1.1 10.9 0.8 52.1 0.5 6.9 0.8 34.0 0.9 52.7 0.4 18.9 0.7 59.6 0.6 63.1 0.3 81.9 0.4 68.5 0.2 91.6 0.1 70.3 0.1 94.0 0.2 4.3 0.5 6.4 0.5 n.d. 15.2 0.1 25.5 0.1 41.3 0.2 60.9 2.5 21.2 0.6 46.0 0.4 16.1 0.2 39.9 0.2 8.2 0.3 3.5 0.1 9.0 0.1 21.5 0.2 36.7 0.2 16.0 0.1 20.3 0.2 18.3 0.1 8.1 0.3 12.2 0.4
29.5 37.5 0.4 0.4 0.4 0.1 0.1 0.1
65.5 53.0 49.3 34.8 30.1 3.9 1.2 0.4
5.0 9.5 50.3 64.8 69.5 96.0 98.7 99.5
Diffusion couple
The primary phases and the solidification sequence were carefully analysed for all alloys. The microstructures of the as-cast alloys show different characteristics, as illustrated in Fig. 1(a)–(h). Of the alloys in the Ti-rich part (> 50 at.% Ti) the as-cast microstructures of alloys 01 and 03 show only the two phases b-Ti and (Fe,Ni)Ti, with b-Ti as the primary phase in alloy 01 and (Fe,Ni)Ti as the primary phase in alloy 03 (Fig. 1(a)). The as-cast microstructures of the alloys 04–07, 16, 32, 33 and 35 show a three-phase microstructure of b-Ti, (Ni,Fe)Ti2 and (Fe,Ni)Ti, with (Fe,Ni)Ti as the primary phase. However, they show two different phase arrangements, i.e. different solidification sequences. In alloys 04, 07 and 32 (Fig. 1(b)) primary dendrites of (Fe,Ni)Ti are embedded within b-Ti and only a few small precipitates of (Ni,Fe)Ti2 are observed. In contrast, alloys 05, 06, 16, 33 and 35 (Fig. 1(c)) show primary (Fe,Ni)Ti surrounded by (Ni,Fe)Ti2 which apparently formed via the peritectic reaction L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2. In between the (Ni,Fe)Ti2 grains some small regions with a lamellar eutectic microstructure composed of b-Ti and (Ni,Fe)Ti2 are found (encircled in black in Fig. 1(c)), which formed from the last liquid. Alloy 15 shows the phases b-Ti and (Ni,Fe)Ti2 in the as-cast state, with b-Ti as the primary phase (Fig. 1(d)).
‘‘Liquid’’ (Fe,Ni)Ti (Fe,Ni)Ti Fe2Ti Fe2Ti a-Fe a-Fe g-Fe n.d., not possible to determine.
Of the alloys in the Ti-lean part (<50 at.% Ti) alloys 11 (Fig. 1(g)), 20 (Fig. 1(e)), 21 show a eutectic microstructure or small regions of eutectic, as in the case of alloy 08 (Fig. 1(f)). These eutectic regions as well the phases involved were analysed by EPMA. The same eutectic composition was observed for alloys 20 and 21, although they show different primary phases. The primary phase in alloy 20 (Fig. 1(e)) is Fe2Ti (29.3Ni; 33.7Ti; 37.0Fe), which is surrounded by a eutectic microstructure of the composition 46.0Ni; 35.8Ti; 18.2Fe. The eutectic consists of the three phases (Fe,Ni)Ti (41.1Ni; 40.0Ti; 18.9Fe), Ni3Ti (51.6Ni; 33.1Ti; 15.3Fe) and Fe2Ti (37.8Ni; 34.4Ti; 27.8Fe).
L.I. Duarte et al. / Intermetallics 18 (2010) 374–384 Table 3 Compositions of coexisting phases measured by EPMA for alloys annealed at 800 C. Alloy no.
Phase
Ni (at.%)
Ti (at.%)
Fe (at.%)
01
b-Ti
0.7 0.1 5.1 0.1 0.6 0.1 4.7 0.1 2.1 0.1 6.3 0.1 4.1 1.1 9.8 0.1 8.0 0.3 15.0 0.1 16.5 0.5 17.7 60.2 30.7 68.2 37.1 72.7 57.4 12.9 58.5 32.1 8.9 0.1 11.1 0.1 3.3 0.1 2.2 0.1 1.0 0.1 0.6 0.1 5.2 0.1 24.9 0.1 17.3 0.2 22.4 0.1 8.1 0.3 3.7 0.3 30.9 17.5 50.6 31.4 18.6 52.0 38.9 0.4 69.2 3.1 10.0 0.5 2.1 0.1 11.6 0.1 15.0 0.3 25.1 0.3
82.3 0.3 67.3 0.2 82.6 0.1 67.2 0.2 51.7 0.3 67.1 0.1 52.0 0.4 67.2 0.1 52.2 0.9 67.4 0.1 52.4 0.4 28.5 22.6 10.7 22.2 5.8 20.8 3.6 27.6 21.6 11.5 29.3 0.2 4.6 0.5 28.7 0.7 3.9 0.7 29.0 0.4 3.2 0.2 91.7 0.1 68.4 0.1 67.5 0.3 52.1 0.3 49.1 0.8 35.9 1.0 47.8 34.5 34.1 46.5 34.1 32.0 49.1 0.3 28.3 2.2 66.7 0.2 86.0 0.2 68.0 0.3 66.7 0.3 66.6 0.3
17.0 0.3 27.6 0.2 16.8 0.1 28.1 0.2 46.2 0.3 26.6 0.1 43.9 0.7 23.0 0.1 39.8 0.6 17.6 0.1 31.1 0.2 53.8 17.2 58.6 9.6 57.1 6.5 39.0 59.5 19.9 56.4 61.8 0.2 84.3 0.5 68.0 0.7 93.9 0.7 70.0 0.4 96.2 0.2 3.1 0.1 6.7 0.1 15.2 0.2 25.5 0.2 42.8 1.1 60.4 1.3 21.3 48.0 15.3 22.1 47.3 16.0 12.0 0.6 2.5 0.9 23.3 0.1 11.9 0.2 20.4 0.4 18.3 0.3 8.3 0.3
03
04 05 06 08
09 10 11
12 13 14 15 16 18 19
20
21 33 34 35 36
(Ni,Fe)Ti2 b-Ti (Ni,Fe)Ti2 (Fe,Ni)Ti (Ni,Fe)Ti2 (Fe,Ni)Ti (Ni,Fe)Ti2 (Fe,Ni)Ti (Ni,Fe)Ti2 (Fe,Ni)Ti Fe2Ti Ni3Ti g-(Fe,Ni) Ni3Ti g-(Fe,Ni) Ni3Ti g-(Fe,Ni) Fe2Ti Ni3Ti g-(Fe,Ni) Fe2Ti g-(Fe,Ni) Fe2Ti a-Fe Fe2Ti a-Fe b-Ti (Ni,Fe)Ti2 (Ni,Fe)Ti2 (Fe,Ni)Ti (Fe,Ni)Ti Fe2Ti (Fe,Ni)Ti Fe2Ti Ni3Ti (Fe,Ni)Ti Fe2Ti Ni3Ti (Fe,Ni)Ti Fe2Ti (Ni,Fe)Ti2 b-Ti (Ni,Fe)Ti2 (Ni,Fe)Ti2 (Ni,Fe)Ti2
When no standard deviation is given for an analysis only few points could be measured indicating that the phase was either too small for a measurement or that only few particles could be detected.
The as-cast microstructures of the alloys 08 (Fig. 1(f)), 09 and 10 all show primary dendrites of g-(Fe,Ni) surrounded by grains of phases Ni3Ti and g-(Fe,Ni). In addition, in alloy 08 small eutectic regions were observed (marked by arrows in Fig. 1(f)). The as-cast microstructure of alloy 11 (Fig. 1(g)) also shows a eutectic microstructure. The composition of this eutectic is 31.5Ni; 20.9Ti; 47.6Fe and it consists of the three phases g-(Fe,Ni) (33.1Ni; 16.4Ti; 50.5Fe); Ni3Ti (35.0Ni; 21.8Ti; 43.2Fe) and Fe2Ti (25.4Ni; 24.3Ti; 50.3Fe). Alloys 13 and 14 both show microstructures where primary a-Fe is surrounded by a lamellar eutectic of a-Fe and Fe2Ti. In alloy 12 (Fig. 1(h)) primary Fe2Ti is surrounded by g-Fe and some more Fe2Ti, while in alloy 18 primary Fe2Ti is embedded in (Fe,Ni)Ti, which apparently formed via the peritectic reaction L þ Fe2Ti 4 (Fe,Ni)Ti. 3.2. Isothermal section at 1000 C After heat treatment at 1000 C the microstructures of all samples were characterized and the phase compositions were determined by EPMA. Equilibria between the phases b-Ti, (Ni,Fe)Ti2
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and (Fe,Ni)Ti as well as between (Ni,Fe)Ti2 and (Fe,Ni)Ti can be observed for the Ti-rich specimens. Fig. 2(a) shows the three-phase equilibrium in alloy 05, whereas the two-phase equilibrium observed in alloy 06 is presented in Fig. 2b. Two different three-phase equilibria were found for the two alloys presented in Fig. 3. Alloy 20 shows the equilibrium between (Fe,Ni)Ti, Fe2Ti and Ni3Ti (Fig. 3(a)). The other three-phase equilibrium between Fe2Ti, Ni3Ti and g-(Fe,Ni) is found in alloy 11 (Fig. 3(b)). In alloys 02 and 15 melt is present at 1000 C, though the annealing time of 10 min was too short to attain equilibrium. The diffusion couple was prepared in order to provide an overview of the phase reactions in the Ti-rich corner of the ternary system at 1000 C. Fig. 4 shows the schematic drawing of this diffusion couple as well as a backscattered electron image from its reaction zone. Strong interdiffusion and phase reactions occurred, resulting in the formation of a series of layers. The total thickness of these layers is about 25–30 mm. The phases (Ni,Fe)Ti2, (Fe,Ni)Ti, Fe2Ti and a-Fe were found in this reaction zone, as indicated in Fig. 4. Several EPMA line scans were performed across the diffusion zone, and from the resulting concentration profiles phase equilibria were determined by extrapolations to the phase boundaries. A tieline obtained from the ‘‘frozen zone’’ and (Ni,Fe)Ti2 is shown in Fig. 5. The other tie-lines follow the binary Fe-Ti phase diagram, being in good agreement with the binary phase compositions. The compositions of the coexisting phases as established by EPMA on samples quenched from 1000 C are shown in Fig. 5, and all compositions are compiled in Table 2. 3.3. Isothermal section at 800 C After heat treatment at 800 C a three-phase microstructure was found in four alloys on the Ti-rich side. Fig. 6(a) presents an example of the three-phase microstructure of alloy 03, showing the equilibrium between b-Ti, (Fe,Ni)Ti and (Ni,Fe)Ti2. For the same temperature an equilibrium between (Ni,Fe)Ti2 and (Fe,Ni)Ti was found in alloys with higher Ni content, for example for the alloys 05, 06, and 16. In Fig. 6(b) the two-phase microstructure of alloy 06 is presented. The equilibrium b-Ti þ (Ni,Fe)Ti2 was found for the alloys 01, 15 and 34. As an example for this equilibrium the microstructure of alloy 15 is shown in Fig. 7(a). Two different three-phase equilibria were found for the alloys with Ti contents <50 at.%. Alloys 19 and 20 show the equilibrium between (Fe,Ni)Ti, Fe2Ti and Ni3Ti. As an example, the microstructure of alloy 19 is presented in Fig. 7(b). The second three-phase equilibrium is between Fe2Ti, Ni3Ti and g-(Fe,Ni). This was observed in the alloys 08 and 11, Fig. 8(a). The alloys 09 and 10 show two-phase equilibria between the Ni3Ti and g-(Fe,Ni) phases. This equilibrium is represented by alloy 09 in Fig. 8b. As can be seen, these samples have a fine lamellar structure that is difficult to analyse by EPMA. The compositions of the coexisting phases as established by EPMA on samples quenched from 800 C are shown in Fig. 9, and all compositions of the phases are compiled in Table 3. 3.4. DTA results As mentioned above, only the data from the first heating curve were considered because of the high reactivity of Ti with the crucible. The yttria coating of the alumina crucible helped to prevent a strong reaction. However, complete suppression of any reaction between the crucible and the melt was not possible for all specimens. The first onset temperatures of melting, which were estimated for each curve from the intersection of the extrapolated baseline with the back-extrapolated peak, are presented in Table 1. These temperatures represent the solidus temperature in all cases.
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Fig. 1. As-cast microstructures of: (a) alloy 03: Fe31.4Ni2.0Ti66.6; (b) alloy 32: Fe28.4Ni5.0Ti66.6; (c) alloy 35: Fe18.3Ni15.0Ti66.7; (d) alloy 15: Fe4.6Ni14.8Ti80.6; (e) alloy 20: Fe22.5Ni43.5Ti34.0; (f) alloy 08: Fe39.0Ni44.4Ti16.6; (g) alloy 11: Fe50.0Ni29.6Ti20.4 and (h) alloy 12: Fe67.8Ni9.4Ti22.8.
Fig. 10 shows the DTA curves recorded during heating of homogenized alloys with a fixed Ti content of 66.7 at.%. The results of these samples are presented in order to obtain a better understanding of the formation of (Ni,Fe)Ti2. It is observed that the solidus temperature decreases monotonously with increasing Ni content to a value of about 1020 C 3 C for the alloys between 10.0 and 15.0 at.% Ni.
For 15.8 at.% Ni a slightly higher onset temperature of 1024 C 3 C is seen, while for the specimen with the highest Ni content of 25.1 at.% an onset temperature of 1018 C 3 C is observed. The onset temperature of 1028 C 3 C found for the alloys with a Ni content of 6.5 at.% suggests that an invariant reaction takes place at about this composition; see Section 4.2.
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Fig. 2. (a) Three-phase microstructure of alloy 05: Fe23.2Ni10.3Ti66.5 and (b) two-phase microstructure of alloy 06: Fe17.7Ni15.8Ti66.5 after annealing at 1000 C for 1167 and 375 h, respectively.
Fig. 3. Three-phase microstructure of two alloys after annealing at 1000 C for 1167 h. (a) Alloy 20: Fe22.5Ni43.5Ti34.0; (b) alloy 11: Fe50.0Ni29.6Ti20.4.
The liquidus temperature could be determined for some alloys by extrapolating the end of the last peak before the return to the baseline, as shown in Fig. 10 for two examples. Unfortunately, in some cases this effect is not well-defined, or the experiment was stopped before melting was completed in order to prevent reactions with the crucible. The determination of the liquidus temperatures is shown in an example in Fig. 10. Liquidus temperatures were only derived from experiments where no reactions with the crucible were observed on the DTA curves and where, by optical inspection after the experiment, no reaction was noticed between sample and crucible. The liquidus temperatures are also given in Table 1.
g-(Fe,Ni) was extensively discussed by Gupta [1], who compared the data with the work done by Drake [16] in which the maximum solubility of Ti in g-(Fe,Ni) was found to be as high as 12–13 at.%. Gupta [1] used the latter data for the boundary of the g-(Fe,Ni) phase, also taking into consideration the work done by Abramycheva et al. [11]. The results of the tie-lines obtained in this work and the resulting boundary for the g-(Fe,Ni) phase, as well as the position of the three-phase field Fe2Ti þ Ni3Ti þ g-(Fe,Ni), are in good agreement with the data assessed by Gupta [1]. It may be speculated that coherency stresses are responsible for the different
4. Discussion 4.1. Isothermal sections at 1000 and 800 C The tie-lines and tie-triangles obtained from evaluation of the individual alloys heat-treated at 1000 C and 800 C are shown in Figs. 5 and 9 and the isothermal sections established from these data are shown in Figs. 11 and 12, respectively. The observation that the compositions of the alloys fall onto the tie-lines or lie within the tie-triangles confirms that the actual compositions are very close to those intended. The results obtained for 1000 C from the alloys with lower Ti content show good agreement with the tie-lines from literature [11,14], as can be seen in Fig. 5. It is noted that some investigators [6,12,15] reported a significantly lower solubility limit of Ti in g-(Fe,Ni). This scattering of the solubility limit of Ti in the phase
Fig. 4. Diffusion couple (schematic illustration and BSE image of reaction zone) between Fe and NiTi2 annealed for 30 min at 1010 C in a vacuum furnace.
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Fig. 5. Isothermal section of the Fe–Ni–Ti system at 1000 C showing EPMA data measured on individual alloys and with experimental data from literature. Empty squares represent the phase compositions according to the accepted assessments of the binary systems.
data reported for the solid solubility of Ti in g-(Fe,Ni), as these stresses may vary considerably in dependence on annealing time and/or cooling rate. According to the present results, the solid solubility for Ti in g-(Fe,Ni) decreases with decreasing temperature (c.f. Figs. 11 and 12) in accordance with the latest assessment [4] and as expected from the data in [7,16]. The maximum solid solubility for Ti in g-(Fe,Ni) is found at the three-phase equilibrium g-(Fe,Ni) þ Fe2Ti þ Ni3Ti. This result also agrees with the experimentally determined isotherms for 900 C [7] and 1027 C [16]. However, the present data contradict earlier data from Speich [6], according to which the solid solubility for Ti in g-(Fe,Ni) increases continuously with increasing Ni content at 1000 C, and data by Abramycheva et al. [11], according to which the solid solubility for Ti in g-(Fe,Ni) at 1000 C has its maximum not at the three-phase equilibrium but at a Ni content which is about 10 at.% higher. Therefore, qualitatively the data confirm that the solubility of Ti in g-(Fe,Ni) varies in dependence on the Fe:Ni ratio according to what was found at 1027 C [16] and at 900 C [7].
Quantitatively, the maximum solid solubility of Ti in g-(Ni,Fe) is given by the results for alloys 11 and 08 at 800 C, which yield 11.5 and 10.7 at.% Ti, respectively (see Table 3), and the maximum value at 1000 C of 12.0 at.% Ti is found for alloy 11. The main differences between this work and the isothermal section at 1000 C proposed by Abramycheva et al. [11] exist in the Ti-rich part, i.e. at a Ti content > 50 at.%. In this part Abramycheva et al. [11] investigated only one diffusion couple between the binary alloys Fe21.7Ni78.3 and Fe43.5Ti56.5 (in at.%). In this diffusion couple they found the phase equilibria (Fe,Ni)Ti þ b-Ti and (Fe,Ni)Ti þ (Ni,Fe)Ti2 from which they deduced the position of the tietriangle (Fe,Ni)Ti þ b-Ti þ (Ni,Fe)Ti2. All other phase equilibria given by Abramycheva et al. [11] in the Ti-rich part are given by dashed lines, indicating that they are tentative until experimental evidence exists. The three-phase equilibrium (Fe,Ni)Ti þ b-Ti þ (Ni,Fe)Ti2 was also found in the present investigation, though a lower Fe content for (Ni,Fe)Ti2 was measured. However the microstructures of several alloys are not in agreement with the isothermal section
Fig. 6. (a) Three-phase microstructure of alloy 03: Fe31.4Ni2.0Ti66.6 and (b) two-phase microstructure alloy 06: Fe17.7Ni15.8Ti66.5 after annealing at 800 C for 1119 and 375 h, respectively.
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Fig. 7. (a) Two-phase microstructure of alloy 15: Fe4.6Ni14.8Ti80.6 and (b) three-phase microstructure of alloy 19: Fe32.8Ni27.2Ti40.0 after annealing at 800 C for 375 and 1100 h, respectively.
Fig. 8. (a) Three-phase microstructure of alloy 08: Fe39.0Ni44.4Ti16.6 and (b) two-phase microstructure alloy 09: Fe29.6Ni55.4Ti15.0 after annealing at 800 C for 1119 h.
proposed in Ref. [11]. For example, according to [11] the alloys 05, 06 and 16 should show a two-phase equilibrium between b-Ti and (Fe,Ni)Ti. Instead, an equilibrium between (Ni,Fe)Ti2 and (Fe,Ni)Ti was found, as can be seen in Fig. 2(b). This means that the phase (Ni,Fe)Ti2 has a much wider homogeneity range at 1000 C than assumed in [11]. In consequence, a second three-phase equilibrium (Fe,Ni)Ti þ bTi þ (Ni,Fe)Ti2 and the three-phase equilibrium (Fe,Ni)Ti þ b-Ti þ L, which were both proposed by Abramycheva et al. [11], do not exist. Instead, the two three-phase equilibria L þ b-Ti þ (Ni,Fe)Ti2 and L þ (Fe,Ni)Ti þ (Ni,Fe)Ti2 exist at 1000 C. Though these three-phase equilibria were not observed in any of the present samples because they include the liquid phase, the position of the respective tietriangles can be localised by the adjoining two-phase equilibria that were experimentally determined. For L þ (Fe,Ni)Ti þ (Ni,Fe)Ti2 the composition of L should be at a slightly higher Fe content than that determined from the diffusion couple (c.f. Fig. 4), and the composition of (Ni,Fe)Ti2 should be at a slightly lower Fe content than that observed for this phase in alloy 36, while the composition of (Fe,Ni)Ti must lie between the compositions of (Fe,Ni)Ti determined in alloy 36 and in the diffusion couple. The phase compositions of the three-phase equilibrium L þ b-Ti þ (Ni,Fe)Ti2 are more difficult to fix. The composition of b-Ti should include the two compositions determined for this phase in alloys 02 and 15, and the composition of (Ni,Fe)Ti2 should be at a slightly higher Fe content than that determined in alloy 02. For the composition of L no data exist, and it can therefore only be estimated. Because both threephase equilibria were not observed directly they are given as dashed lines in Fig. 11.
So far no isothermal section at 800 C is available in literature. The isothermal section in Fig. 12 was based on the current EPMA results shown in Fig. 9. Even after prolonged annealing at this
Fig. 9. Compositions of alloys (solid black circles) and compositions of phases as determined by EPMA (empty black circles) of the Fe–Ni–Ti system at 800 C.
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Fig. 10. DTA curves of the Ti-rich alloys with 66.7 at.% of Ti obtained during heating at 10 K/min. Onset temperatures of melting (left) and liquidus temperatures (right) are indicated.
temperature, it is difficult to generate phases that are large enough for EPMA measurements, as can be seen in Figs. 3(b) and 8 for the Ni3Ti phase. Therefore, concentrations in the Ti-lean part of the system are somewhat uncertain and are indicated by dashed lines. The fine lamellar structure observed in Fig. 8 for both alloys is due to the ‘‘cellular precipitation reaction’’ described by Speich [6] and Fournelle [12]. According to these reports the decomposition of Fe-30Ni-6Ti (in wt.%) occurs via a cellular precipitation reaction in which cells consisting of alternate lamellae of Ni3Ti (DO24) and austenite are formed. This first cellular reaction is subsequently consumed by a second cellular reaction of much coarser spacing. These authors also reported that Ni3Ti forms an almost perfectly coherent interface with the austenite, with the following orientation relationship:
gB þ Ni3Ti / gC þ (Fe,Ni)2Ti, where gA, gB and gC are austenitic solid solutions of different solute content [6]. As well as in Fig. 8 these features can also be observed in Fig. 3(b), where lamellar Ni3Ti grew inside of the g-(Fe,Ni) grains. No ternary compound could be found at both temperatures (Figs. 11 and 12). Compared to the section at 900 C by van Loo et al. [7], the same phase equilibria were observed at 800 C, with all binary compounds showing a large solubility for the third element. From 900 C to 1000 C, phase equilibria in the Ti-rich corner change because above 942 C a melt forms within the Ni–Ti binary system.
ð001ÞNi3 Ti ==ð111Þg; ½010Ni3 Ti ==½110g
From the examination of the as-cast microstructures and the DTA investigations the liquidus projection shown in Fig. 13 was established. In the Ti-rich part, (Fe,Ni)Ti was observed as a primary phase in all alloys with 66.7 at.% Ti which contained between 2 and 25.1 at.% Ni (alloys 03–06, 32, 33, 35, 36). For the alloys with up to 5.0 at.% Ni (alloys 03, 04, 32), primary dendrites of (Fe,Ni)Ti are surrounded by b-Ti (Fig. 1(a and b)) although only in the case of alloy 03 is there no indication of any reaction between these two phases. In addition, alloys 04 and 32, i.e. those with the highest Ni content, show small amounts of eutectic (Ni,Fe)Ti2 þ b-Ti (Fig. 1(b)) which should be formed from the last solidifying melt. Apparently the (Ni,Fe)Ti2 that surrounds the (Fe,Ni)Ti, (Fig. 1(b)) is already formed in alloys 04 and 32 by the peritectic reaction L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2. The latter feature was also observed by [18]. In contrast, in the alloys with Ni content of between 10.0 and 15.8 at.% Ni (05, 06, 35, 33) (Fe,Ni)Ti is surrounded by (Ni,Fe)Ti2, with clear evidence that (Fe,Ni)Ti have partly dissolved in order to form (Ni,Fe)Ti2 (Fig. 1(c)), i.e. (Ni,Fe)Ti2 formed via the peritectic reaction L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2. The above observations and also those made for the other alloys in the Ti-rich part are in full agreement with the liquidus projection proposed by Riani et al. [18], and the Ti-rich part in Fig. 13 was drawn accordingly. In addition, the temperatures for the invariant reaction U2: L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 þ b-Ti (at about 6 at.% Ni) of about 1030 C and the maximum (M2) on the monovariant line L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 at approximately 1035 C established here agree well with the data assessed by Cacciamani et al. [3], who gave w1025 C and w1030 C, respectively. It is noted that a recent thermodynamic modelling of this system published by
It was reported by Speich [6] that the ternary alloys may exhibit two types of discontinuous precipitation: 1) gA / gB þ Ni3Ti and 2)
Fig. 11. Isothermal section of the Fe–Ni–Ti system at 1000 C. Empty squares represent the phase compositions according to the accepted assessments of the binary systems.
4.2. Liquidus surface and reaction scheme
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Fig. 12. Isothermal section of the Fe–Ni–Ti system at 800 C. Empty squares represent the phase compositions according to the accepted assessments of the binary systems.
De Keyzer et al. [19] yields slightly higher temperatures for these invariant reactions (U2 ¼ 1043 C and M2 ¼ 1044 C). In contrast to the Ti-rich part the solidification behaviour of the Ti-lean alloys was found to differ from liquidus projections established previously, and therefore this part will be discussed in more detail. The solidification of alloys 13 and 14 both end at the eutectic trough e2–U1, and the Ni-rich alloy (13) at a somewhat lower temperature, which is in full agreement with the run of the eutectic trough, which descends from 1289 C (e2) in the Fe–Ti binary system to about 1199 C at the invariant reaction U1. The temperature for U1 was not observed directly but is estimated from the onset temperature observed for alloy 12, which is 1199 C (Table 1). From the composition of alloy 12 and the composition of the primary Fe2Ti phase (7.9Ni; 27.8Ti; 64.3Fe) in that alloy we conclude that the effect at 1199 C belongs to the reaction U1. This temperature is in good agreement with that of 1200 C given in the last assessments [3,4], although the calculated value (1270 C) [19] is much higher. While onset temperatures and metallographic observations for alloys 13 and 18 confirm that the peritectic reaction L þ Fe2Ti 4 (Fe,Ni)Ti continues from p2 with decreasing temperature into the ternary system, observations for alloys 20 and 21 contradict previous liquidus projections [1,3,4,18]. According to these previous works the solidification of alloys 20 and 21 will end up at the eutectic trough e3–U2 (labelled e3–E2 in Fig. 13; see below), whose temperature was supposed to descend towards to the binary system from U2 1118 < T < 1317 C to e3 1118 C. Since the first onsets for both alloys are below 1118 C and because the Ni-rich alloy (21) has the higher onset (1113 C), this means that the eutectic trough descends into the ternary and not towards the binary. Moreover, the onset for alloy 20 is at 1099 C, which is the lowest temperature observed for all the Ti-lean alloys. That the eutectic trough descends inward and that the lowest onset temperature is observed here consequently means that the invariant reaction U2 is actually a ternary eutectic (E2: L 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti). This conclusion is also confirmed by the EPMA results which show that a eutectic is located at 46.0Ni; 35.8Ti; 18.2Fe and that it consists of the phases Fe2Ti, Ni3Ti and (Fe,Ni)Ti.
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Fig. 13. Proposed liquidus projection of the Fe–Ni–Ti system according to this study.
In view of the large volume fraction of this ternary eutectic (Fig. 1(e)) the onset temperature of 1099 C for alloy 20 should actually be the temperature of E2. This temperature is in good agreement with that recently calculated by De Keyzer et al. (E2 ¼ 1103 C) [19], although the calculated composition is at a much higher Ni content than that established by EPMA. The DTA results for alloys 08–10 (Table 1) show that the onset temperatures decrease from alloy 10 to 08, agreeing with previous work. The solidification of alloys 09 and 10 ends up at the eutectic trough e1–E1, with the Ni-rich alloy (10) showing a somewhat higher solidification temperature (1267 C). This agrees fully with the run of the eutectic trough, which descends from 1304 C from the binary eutectic e1 towards E1. According to previous assessments [1,3,4,18] and the calculated liquidus projection [19] alloy 08 should end up at the ternary eutectic E1 at a temperature of around 1110 C and be the closest alloy to this reaction. In fact, the microstructure of the as-cast alloy 08 shows only small amounts of eutectic regions (Fig. 1(f)). That the signal observed in the DTA for E1 in alloy 08 (at 1109 C) is only small is consistent with this observation. Actually, EPMA investigations on alloy 11 revealed that the composition of the eutectic E1 (31.5Ni; 20.9Ti; 47.6Fe) is close to this alloy, i.e. at a considerable lower Ni content than previously assumed. According to the first onset observed for alloy 11 in the DTA the temperature for E1 is 1108 C. Between the two ternary eutectics there must be a maximum in temperature on the invariant line L þ Fe2Ti þ Ni3Ti. According to Alkemade’s theorem [20] this maximum should be at the point where the connecting line between the congruent compositions of Fe2Ti and Ni3Ti (so-called Alkemade line) intersects with the invariant line. This point is represented by M1 in Fig. 13 and its temperature is estimated from the course of the isotherms to be at about 1150 C, which is a little bit lower than the temperature of 1187 C calculated by De Keyzer et al. [19]. In view of the above, the liquidus projection as well as the reaction scheme must be modified as follows: a) U2 becomes E2: L 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti at about 1099 C; b) E1: L 4 g(Fe,Ni) þ Fe2Ti þ Ni3Ti at 1108 C; c) the line E1–E2 has a maximum (M1) at around 1150 C. It is noted that the modified liquidus
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Fig. 14. Reaction scheme of the Fe–Ni–Ti system.
projection (Fig. 13) and the reaction scheme (Fig. 14) are consistent with one another. 5. Conclusions Isothermal sections of the ternary system Fe–Ni–Ti were established at 1000 C and for the first time at 800 C. A good correlation was found between our results and those from literature for both isothermal sections. Qualitatively the present data confirm that the solid solubility of Ti in g-(Fe,Ni) a) varies in dependence on the Fe:Ni ratio, as found before at 1027 C [16] and at 900 C [7]; and b) decreases with decreasing temperature (c.f. Figs. 11 and 12). The liquidus projection as well as the reaction scheme in the Tilean part must be modified as follows: a) U2 becomes E2: at 1099 C; b) E1: L 4 gL 4 Fe2Ti þ Ni3Ti þ (Fe,Ni)Ti (Fe,Ni) þ Fe2Ti þ Ni3Ti at 1108 C; c) the line E1–E2 has a maximum (M2) at around 1150 C. The invariant reaction in the Ti-rich part appears to be of quasiperitectic four-phase transition-type (U) L þ (Fe,Ni)Ti 4 (Ni,Fe)Ti2 þ bTi at around 1030 C. This implies the existence of a maximum (M1) on the monovariant liquidus line, at approximately 1035 C, as already suggested by Cacciamani et al. [3] and Riani et al. [18]. Further, a three-phase equilibrium between b-Ti, (Fe,Ni)Ti and (Ni,Fe)Ti2 was observed in some alloys, but neither the three-phase equilibrium b-Ti þ L þ (Fe,Ni)Ti, nor the two-phase equilibrium bTi þ (Fe,Ni)Ti were found in any sample. Instead, an equilibrium between (Ni,Fe)Ti2 and (Fe,Ni)Ti was observed. Alternative liquidus projection and reaction schemes were proposed based on these new experimental data.
Acknowledgements This work was performed within the framework of the COST Action 535, ‘‘Thermodynamics of alloyed aluminides (THALU)’’, and was financially supported by the Switzerland State Secretariat for Education and Research (SER) under Contract no. C05.0031.
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