On the reaction scheme and liquidus surface in the ternary system Fe–Si–Ti

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Intermetallics 16 (2008) 273e282 www.elsevier.com/locate/intermet

On the reaction scheme and liquidus surface in the ternary system FeeSieTi Franz Weitzer a, Julius C. Schuster a,*, Masaaki Naka b, Frank Stein c, Martin Palm c b

a Innovative Materials Group, Universita¨t Wien, Wa¨hringer Straße 42, A-1090 Wien, Austria Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki, Osaka 567, Japan c Max Planck Institut fu¨r Eisenforschung GmbH, Max Planck Straße 1, D-40237 Du¨sseldorf, Germany

Received 28 June 2007; accepted 11 October 2007 Available online 26 November 2007

Abstract The constitution between 900  C and 1700  C of the ternary system FeeSieTi is investigated over the entire composition range by X-ray diffraction (XRD), differential thermal analysis (DTA), electron probe microanalysis (EPMA), and scanning electron microscopy/energy dispersive X-ray (SEM/EDX). Nine ternary phases are observed. The crystal structures for the phases t1-FeSi2Ti and t2-FeSiTi are confirmed. The crystal structure of the phase t3-Fe4Si3Ti is shown to be of Pd40Sn31Y13-type (space group P6/mmm, a ¼ 1.72073(8) nm and c ¼ 0.79819(6) nm). Four more ternary phases, t5, t7, t8, and t9, are stable at 900  C. Their compositions were determined, but their crystal structures are still under investigation. The tie lines observed for this temperature agree with the literature data reported for 800  C and 1000  C, respectively. The phases t4 and t6 occur at higher temperatures and are not equilibrium phases at 900  C. The liquidus surface shows wide regions of primary solidification for the congruently melting phases t1 and t2, small regions of primary solidification for the phases t3et6, t8 and t9, but none for t7. A reaction scheme linking the liquidus surface with the equilibria at 900  C is proposed. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Ternary alloy systems; B. Phase diagrams

1. Introduction For ceramic joining applications filler metals are needed, which have low melting temperatures and still sufficient reactivity. Silicon is a typical additive to reduce the melting point of base metals like iron, cobalt or nickel. To enhance the reactivity of filler metals further additions such as Ti are commonly used. Previous research showed such NieSieTi alloys to be promising candidate materials [1]. These results motivated us to study FeeSieTi filler metals, too. Thus, in order to understand the build up of the interfacial reaction zone an investigation of the phase equilibria of the Si-rich part of the system FeeSieTi became desirable, because previous literature reports only on the iron-rich section of this system.

* Corresponding author. E-mail address: [email protected] (J.C. Schuster). 0966-9795/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2007.10.006

Such interest in the constitution and an ensuing thermodynamic description of the entire FeeSieTi system was reinforced by recent efforts to obtain metallic FeeTi alloys by direct reduction of ilmenite (FeTiO3) in liquid silicon [2]. Vogel and Schlu¨ter were the first to explore the iron-rich section up to FeSieFeSiTieFe2Ti [3]. Using thermal analysis and LOM-metallography they reported a liquidus projection showing entirely ternary primary crystallization regions for the phase Fe5Si3 (then ‘‘Fe3Si2’’) containing Si, as well as for the ternary phase FeSiTi. The phase ‘‘Fe3Si2’’ was found to form from L þ FeSi þ FeSiTi at 1200  C. The composition (converted into at%) of the liquid phase was given (approximately) as Fe63Si29.5Ti7.5. A ternary eutectic L ¼ a2-Fe3Si þ ‘‘Fe3Si2’’ þ FeSiTi occurred at 1145  C and (approximately) Fe55Si37Ti8. The eutectic trough between a,a2(Fe) and Fe2Ti shows a maximum near 5 at% Si. Furthermore, three transition reactions and a peculiar, more Fe-rich eutectic were reported. Extended solubility of Si in the

F. Weitzer et al. / Intermetallics 16 (2008) 273e282

274

Table 1 Compositions and crystal structures of the ternary phases in FeeSieTi Ternary phase

EDX data (at%) Fe

Si

Reference

t2-FeSiTi t3-Fe40Si31Ti13

t4 t5 t6 t7 t8 t9 Fe7Si2Ti Fe3xSiTix

49e50 33 33e35 35.5 36 36 45.6 64.3 w49 40 40 43 20

Structure type

Lattice parameters (nm) a

b

c

Pbam

MnSi2Ti

Ima2

FeSiTi

0.9534(1) 0.95427(7) 1.0830 1.0827(1)

P6/mmm

Pd40Sn31Y13

0.86137(8) 0.86115(6) 0.6997 0.69869(7) 1.72073(8) 1.70789 1.7206

0.76396(4) 0.76313(6) 0.6287 0.62991(6) 0.79819(6) 0.7971 0.7981

I-43m Fm3m

Fe5Si2V3 MnCu2Al

Ti

t1-FeSi2Ti 24e25 30e35 31e33 49 49 52 28 7.4 w12.5 10 20 17 69

Space group

25e26 32e37 33e35 15.5 15 12 26.3 28.3 w38.5 50 40 40 11

This work [9] This work This work [9] [7,8] This work This work This work This work This work This work [9]

Laves-phase Fe2Ti was noted. According to Yan et al. [4] Si is replacing iron equally on both crystallographic positions 2a and 6h suggesting a formula (Fe1xSix)2Ti. The melting temperatures for such Laves-type solid solution with the composition Fe56Si19Ti25 as well as for the ternary phase FeSiTi were determined to be 1470  C [5] and 1760  20  C [6], respectively. Solid state phase equilibria over the entire FeeSieTi composition range were reported by Markiv et al. [7] (reproduced in Ref. [8]). Using XRD and metallography, these authors observed five ternary phases in the isothermal section at 800  C: t1-FeSi2Ti, t2-FeSiTi, t3 (at Fe52Si36Ti12, corresponding to ‘‘Fe4Si3Ti’’), X0 (at Fe10Si44Ti46), and X00 (at Fe15Si40Ti45). The extensive solid solubility of Si (substituting for Fe) in the Laves-phase Fe2Ti was corroborated. Very recently Lo¨ffler [9] reinvestigated the equilibria for <40 at% Si by EPMA and XRD, and presented isothermal sections at 800  C, 1000  C, and 1150  C (reproduced in

0.8837 0.5709

Reference

[20] This work [24] This work This work [9] [25]

[9] [15]

part in Refs. [10,11]) as well as a table with DTA data, which were interpreted only partially. This work confirmed the large Si-solubility in Fe2Ti, the stabilisation of Fe5Si3 by Ti towards lower as well as higher temperatures (relative to the binary stability range), and the existence of t2-FeSiTi at all temperatures was investigated. The phase t3-‘‘Fe4Si3Ti’’ was found at 1150  C and 1000  C, but contrary to Markiv et al. [7] not at 800  C. The typical composition was Fe49Si36Ti15. An additional cubic Fe-rich phase, Fe7Si2Ti, was seen at 1150  C and 1000  C, but not at 800  C. The monovariant eutectic reaction L ¼ a,a2(Fe) þ Fe2Ti (with Si in solid solution) was observed at 1317  C (in alloy Fe75.4Si6.4Ti18.2) and at 1311  C (in alloy Fe69.1Si12.7Ti18.2). Both temperatures are higher than the binary value of 1293  C, confirming thus the maximum in the eutectic trough reported by Vogel and Schlu¨ter [3]. A metastable Heusler-type phase Fe2SiTi having a lattice parameter exactly twice that of a-Fe and promoting hardening

Fig. 1. XRD powder pattern for t3 (Cu Ka1-radiation) observed in alloy Fe49Si35.5Ti15.5 equilibrated at 900  C. Indexing of the pattern is based on a hexagonal unit cell with a ¼ 1.72073(8) nm and c ¼ 0.79819(6) nm.

F. Weitzer et al. / Intermetallics 16 (2008) 273e282

of ferrite was repeatedly reported [10,12e16] to form at or below 800  C. The g-Fe region was found to have a very limited extension in FeeSieTi [17]. The early proposals for the crystal structures of t1-FeSi2Ti and t2-FeSiTi by Markiv et al. [7,18,19] were not corroborated later. Steinmetz et al. [20] reported FeSi2Ti to be of MnSi2Titype (space group Pbam, oP48). In an apparently independent structure determination Yarmolyuk et al. [21] found FeSi2Ti to be of CrSi2Zr-type (also space group Pbam, oP48). Crystal structure standardisation [22] shows both structure proposals to be identical. The crystal structure of FeSiTi was determined by Jeitschko [23,24] to be a new structure type (space group Ima2, oI36). Lattice parameters were reported for t3-‘‘Fe4Si3Ti’’ by Steinmetz et al. [25] and confirmed by Lo¨ffler [9]. Last not least, the latter author determined the Fe-rich phase Fe7Si2Ti to have Fe5V3Si2-type crystal structure. The present work was initiated with the aim to establish the solidification sequence in the Si-rich region. It was possible to extend this investigation over a much wider composition and temperature range, because the solid state data collected in the thesis of Lo¨ffler [9] became available very timely. Throughout the present study the phase diagrams of the binary boundary systems FeeTi, FeeSi, and TieSi were accepted as assessed in the book of Massalski [26]. The only amendment to be made regards the phase(s) Ti5Si4, because contrary to the TieSi phase diagram assessment [26], this phase exists in two modifications as found by Nickl and Sprenger [27] in samples prepared by vapor deposition: a tetragonal modification (space group P412121, Zr5Si3-type a ¼ 0.6702 nm and c ¼ 1.274 nm), which did not decompose upon annealing at 850  C, and an orthorhombic modification (space group Pnma, a ¼ 0.66645 nm, b ¼ 0.6505 nm, c ¼ 1.2690 nm), which did decompose into the tetragonal phase upon annealing at 1100  C. Based on these

Fig. 2. Isothermal section for 900  C (note: phase boundaries shown are schematic only). The phase fields of a(Fe) and a2-Fe3Si are not separated and the term a,a2(Fe) is used.

275

observations the tetragonal modification was assumed by Nickl and Sprenger to be a low temperature modification, though it has the higher symmetry. 2. Experimental More than 80 ternary FeeSieTi alloys were synthesized by arc melting under argon (from ingots of 99.98 wt% pure Fe and Ti, as well as 99.99 wt% pure Si; all supplied by Johnson Matthey Alfa Products, Karlsruhe, Germany). A piece of each alloy was inspected in the as cast state, while other parts were heat treated in evacuated quartz tubes for a minimum of two weeks at 900  C, 950  C, or 1000  C followed by water quenching. Phase identification was done by X-ray powder diffraction using image foil equipped Guinier type chambers (model 670, Huber Diffraktionstechnik, Rimsting, Germany), Cu Ka1 radiation, 99.9999 wt% pure Ge as internal standard, and for data handling and treatment using the software packages CSD [28] and STRUKTUR [29]. Energy dispersive X-ray analysis (EDX) was done on two different scanning electron microscopes (annealed alloys: model DSM 962, Zeiss Table 2 Comparison of observed and calculated intensities for the XRD peaks of t3Fe40Si31Ti13 shown in Fig. 1 hkl

2qobs.

2qcalc.

Iobs.(%)

Icalc.(%)

3 2 3 4 5 2 4 2 4 5 5 3 4 5 6 5 6 4 4 7 3 4 4 5 7 5 0 5 1 4 5 4 5 7 5

26.04 27.39 28.38 29.69 29.96 30.54 35.59 35.84 36.66 37.65

26.04 27.37 28.38 29.64 29.96 30.53 35.58 35.81 36.65 37.64 37.67 38.63 39.12 39.38 39.63 40.55 41.27 41.68 41.97 42.43 43.04 43.30 43.93 43.98 43.98 44.18 45.41 45.67 46.69 46.94 47.62 48.01 48.19 48.42 49.05

4.5 3.6 10.3 3.3 5.0 21.7 4.8 4.8 8.6 7.5

1.6 4.2 12.8 3.0 7.7 18.5 5.9 8.7 6.5 3.0 4.0 7.4 9.7 3.4 7.0 4.1 9.2 53.8 52.7 5.8 3.6 100 8.8 9.8 38.8 44.5 44.5 8.9 6.4 11.7 12.2 53.4 10.6 10.4 4.0

2 1 2 1 0 2 1 0 3 0 2 3 2 2 1 1 1 0 4 0 2 3 1 3 0 2 0 0 1 2 4 4 1 0 4

0 2 1 1 0 2 2 3 0 2 0 2 2 1 0 2 1 3 0 0 3 2 3 1 1 2 4 3 4 3 0 2 3 2 1

38.65 39.1 39.33 39.63 40.55 41.27 41.68 41.97 42.44 43.08 43.30 43.99

44.19 45.42 45.67 46.70 46.97 47.61 48.02 48.19 48.42 49.06

23.6 8.0 2.9 6.9 4.7 8.7 81.8 57.2 6.6 1.4 100 62.0

49.6 53.0 9.8 10.1 6.6 13.8 53.4 14.3 13.9 2.9

For the calculation the structural data given in Table 1 were used.

276

Table 3 Results of EDX, XRD, and thermal characterization of representative alloys in FeeSieTi Fe (at%)

Si (at%)

Ti (at%)

Heat treatment

Phases present

Space group, prototype, Strukturbericht type

a

b

c

5

47.5

47.5

1000  C

TiSi Ti5Si4(l ) t1-FeSi2Ti

Pnma, FeB, B27 Pnma, Sm5Ge4 Pbam, MnSi2Ti

0.65399(4) 0.64734(4) 0.86125(6)

0.36389(3) 1.26689(8) 0.95271(5)

0.49985(4) 0.66244(4) 0.76331(6)



Lattice parameters (nm)

Si (at%)

Ti (at%)

DTA signals ( C) on heating at a rate of 5 K/min

0

49.9

50.1

1.1

44.0

54.9

1391 1450 1484

U6 U5 P4

Fe (at%)

14

53

33

900 C

TiSi2 TiSi t1-FeSi2Ti

Fddd, TiSi2, C54 Pnma, FeB, B27 Pbam, MnSi2Ti

0.47988(4) 0.65398(4) 0.86147(7)

0.82673(8) 0.36399(3) 0.95384(8)

0.85512(7) 0.50013(4) 0.76334(7)

0 0 24.2

66.6 49.6 50.0

33.4 50.4 25.8

1379 1399 1422

U7 U6 Offset

8

63

29

900  C

TiSi2 t5 t1-FeSi2Ti

Fddd, TiSi2, C54

0.82671(8)

0.47997(5)

0.85513(9)

0.86115(6)

0.95427(7)

0.76313(6)

66.6 64.3 49.9

33.4 28.5 25.4

1263 1372 1421

P6

Pbam, MnSi2Ti

0 7.2 24.7

(Si) t5 t1-FeSi2Ti

Fd-3m, Cdiamond, A4

0.54317(2)

Pbam, MnSi2Ti

0.86170(7)

0.76345(7)

0 7.2 24.9

100 64.4 50.0

0 28.4 25.1

(1174) 1208 1230 1271 1347

Trace U10 U9 Offset Offset

(Si) t5 t1-FeSi2Ti

Fd-3m, Cdiamond, A4

0.54317(2)

Pbam, MnSi2Ti

0.86141(5)

0.95336(6)

0.76295(5)

0 7.4 24.8

100 64.3 50.1

0 28.3 25.1

1208 1268 1345 1362

U10 P6 Offset Offset

(Si) Fe2Si5 t1-FeSi2Ti

Fd-3m, Cdiamond, A4 P4/mmm, Fe2Si5 Pbam, MnSi2Ti

0.54316(2) 0.26910(4) 0.86084(6)

0.95331(7)

0.51435(9) 0.76326(6)

0 32.4 25.1

100 66.2 50.1

0 1.4 24.8

1175 1202 1325

E2 Offset Offset

Fe2Si5 FeSi2 t1-FeSi2Ti

P4/mmm, Fe2Si5 Cmca, FeSi2 Pbam, MnSi2Ti

0.26979(3) 0.98794(8) Trace

0.77997(6)

0.51434(9) 0.78376(7)

30.7 33.2

68.8 66.5

0.5 0.3

1006 1185

Small U14

FeSi2 FeSi t1-FeSi2Ti

Cmca, FeSi2 P213, FeSi, B20 Pbam, MnSi2Ti

0.98905(9) 0.44841(2) 0.86101(6)

0.78029(7)

0.78384(7)

0.95272(6)

0.76342(6)

33.2 50.0 25.6

66.8 50.0 50.0

0 0 24.4

1011 1184 1307 1348

Small U14 Offset Offset

49.4 25.1

50.6 50.0

0 24.9

(1182) 1328

Trace emax3

(1298) 1326 1407

Trace, U8 emax3 Offset

1254 1433

E1 P5

(1253) 1428

Trace P5

10

19

31

36

55

69

65

54

18

24

12

4

10

900 C

900  C

900  C



900 C

900  C

44

45

11

900  C

FeSi t1-FeSi2Ti

P213, FeSi, B20 Pbam, MnSi2Ti

42.5

42.5

15

900  C

FeSi t1-FeSi2Ti

P213, FeSi, B20 Pbam, MnSi2Ti

33.5

41

25.5

1000  C

FeSi t1-FeSi2Ti t2-FeSiTi

P213, FeSi, B20 Pbam, MnSi2Ti Ima2, FeSiTi

0.44859(5) 0.86024(6) 0.70191(9)

FeSi t4 t1-FeSi2Ti t2-FeSiTi

P213, FeSi, B20

0.44893(4)

Pbam, MnSi2Ti Ima2, FeSiTi

0.86032(6) 0.70093(6)

32

41

27

950  C

0.95358(7)

0.94382(7) 1.07955(9)

0.95203(6) 1.08029(9)

0.76451(6) 0.62701(6)

0.76384(5) 0.62746(6)

Offset

F. Weitzer et al. / Intermetallics 16 (2008) 273e282

21

72



31

41

47.9

50.3

60

41

45

34.7

46.0

32

28

14

17.4

3.7

8

900  C

900  C

1000  C

900  C

900  C

Pbam, MnSi2Ti Ima2, FeSiTi P213, FeSi, B20

Trace 0.86048(5) 0.69869(7)

FeSi t4 t2-FeSiTi

Ima2, FeSiTi

0.44876(1) Trace 0.69883(6)

FeSi t2-FeSiTi t3-Fe4Si3Ti

P213, FeSi, B20 Ima2, FeSiTi P6/mmm, Pd40Sn31Y13

0.44879(3) 0.69836(7) 1.72113(6)

FeSi Fe5Si3 t3-Fe4Si3Ti

P213, FeSi, B20 P63/mcm, Mn5Si3, D88 P6/mmm, Pd40Sn31Y13

0.44883(6) 0.67763(7) Traces

Fe5Si3 a2-Fe3Si t3-Fe4Si3Ti

P63/mcm, Mn5Si3, D88 Fm-3m, BiF3, D03 P6/mmm, Pd40Sn31Y13

0.67968(7) 0.56529(8) 1.72100(7)

0.95239(2) 1.0827(1)

0.76526(6) 0.62991(6)

1.0803(1)

0.62664(7)

1.0861(1)

0.62541(6) 0.79715(5)

28.1 25.1 33.6

45.6 48.9 33.5

26.3 26.0 32.9

1414 1480

P5 Offset

50 31

49 45

1 24

(1254) 1259 1304 1322

Shoulder emax5 Offset Offset

1241 1258

P7 emax5

49.5 53.3 49.0

50.5 37.2 36.0

0 9.5 15.0

1199 1251 1365

U12 Offset Offset

36.5 26.7 35.7

7 3.3 15.4

1151 1180

E3 Offset

0.79836(5)

56.5 70 48.9

0.79805(5)

58.4 69.4 49.0

36.6 27.2 35.7

5 3.4 15.3

1151 1194 1201

E3 U13 P8

1.08078(9)

0.62474(6) 0.79801(5)

69 37 49

26 32 35

5 31 16

1160 1184

U16 Offset

26.3 32.9

5.5 33.2

U10

0.52797(5)

68.2 33.9

1198

1.08094(9)

47

16

37

1171 1214

U15 ?

0.47152(8) 0.47268(9)

57

35

8

900 C

Fe5Si3 a2-Fe3Si t3-Fe4Si3Ti

P63/mcm, Mn5Si3, D88 Fm-3m, BiF3, D03 P6/mmm, Pd40Sn31Y13

Traces 0.56498(8) 1.72106(7)

53

31

16

900  C

a2-Fe3Si t2-FeSiTi t3-Fe4Si3Ti

Fm-3m, BiF3, D03 Ima2, FeSiTi P6/mmm, Pd40Sn31Y13

0.56545(9) 0.69675(8) 1.72099(8)

a2-Fe3Si t2-FeSiTi

Fm-3m, BiF3, D03 Ima2, FeSiTi

0.56631(7) 0.69975(6)

Ti5Si3 Fe2Ti FeTi

P63/mcm, Mn5Si3, D88 P63/mmc, MgZn2,C14 Pm-3m, CsCl, B2

Traces 0.48236(4) Traces

Ti5Si3 FeTi (Ti) Ti5Si3 Ti3Si (Ti)

P63/mcm, Mn5Si3, D88 Pm-3m, CsCl, B2 P63/mmc, Mg, A2 P63/mcm, Mn5Si3, D88 P42/n, Ti3P P63/mmc, Mg, A2

0.74495(5) 0.29843(3) Traces 0.74531(2) Traces Traces

0.51490(6)

2 47.5

35 1.5

63 51

1034 1179

E4 U15

0.51459(3)

0.2 0 4.2

34.7 24.9 1.7

65.1 75.1 94.1

1201

Small

Ti5Si3 Fe2Ti t2-FeSiTi

P63/mcm, Mn5Si3, D88 P63/mmc, MgZn2,C14 Ima2, FeSiTi

0.74102(5) 0.48204(4) 0.69866(5)

4.1

37.0

58.9

1591 1637

P3 Offset

1.0946(1)

0.51021(9) 0.77204(9) 0.63208(4)

30.5

37.5

t8 t1-FeSi2Ti t2-FeSiTi

Pbam, MnSi2Ti Ima2, FeSiTi

0.86050(9) 0.70184(7)

0.95264(9) 1.0800(1)

0.76367(7) 0.62769(5)

(1416) 1521 1532 1615

Trace U2 emax2 Offset

1640 1663

P1 Offset

44 45

31 15

25 40

900  C 900  C

29

14

57

900  C

2

23

75

950  C

27.1

31.4

41.5

900  C

23

23

40

36

37

41



1000 C

900  C

t7 t8 t2-FeSiTi

Ima2, FeSiTi

0.70069(5)

0.7812(1)

1.08905(9)

0.62893(5)

32

10.6 19.9 31.7

39.7 38.0 32.9

49.7 42.1 35.4

F. Weitzer et al. / Intermetallics 16 (2008) 273e282



t4 t1-FeSi2Ti t2-FeSiTi

(continued on next page) 277

(1479) 1539 1568 1599

U3 40.4 26.0 42.3 49.4

1485

54.0 49.7 39.1 43.9 41.3 42.8

Trace U1 P2 Offset

55.4 61.9 49.8

0.66254(4) Traces Pbam, MnSi2Ti

1.26810(9) Ti5Si4(l ) t9 t1-FeSi2Ti

36.6

43.8

45

43.9 12.3

18.4

46 42 12

2.5

57

1000 C



0.64747(4)

0.95322(6) 0.86124(5) Pbam, MnSi2Ti t8 t9 t1-FeSi2Ti 900  C

Pnma, Sm5Ge4

0.76396(5)

17.3 24.6

2.1? 9.0 18.1 0.66260(4) 1.26715(9) 0.64652(4) Ti5Si4(l ) t7 t9 1000  C

Pnma, Sm5Ge4

0.66270(5) 0.51391(7) 1.26759(8) 0.64892(5) 0.74463(4) Pnma, Sm5Ge4 P63/mcm, Mn5Si3, D88 Ti5Si4(l ) Ti5Si3 t7 1000 C

0.63145(5) 1.0898(1) 0.70001(9)

40.5



Ima2, FeSiTi

0 1.2 10.2

44.6 36.9 40.0

(1480) 1567 1597 1608

No signal

<1500

emax1 Offset 1662 1675 58.8 48.8 35.1 0.74531(2)

a

Ti5Si3 t7 t2-FeSiTi

The identification of the ternary phases was done primarily by XRD powder diffraction and corroborated by EDX

15

48

The crystal structures reported for t1-FeSi2Ti and t2-FeSiTi (Table 1) are corroborated by comparing experimentally observed XRD powder patterns with the patterns calculated from the structural data. The phase t3-Fe4Si3Ti1 is newly determined to be of Pd40Sn31Y13-type crystal structure [30], isostructural with Ni40Si31Ti13 [31]. The indexed XRD powder pattern is shown in Fig. 2. A comparison between observed and calculated intensities is made in Table 2. For several other ternary phases crystal structure determination is in progress. Thus, at present these ternary phases are labeled only with the greek letter t and an index number. All these phases were observed metallographically, too, and their compositions were determined by EDX (Table 1). The phases Fe7Si2Ti [9] and the supposedly Heusler-type ordered phase Fe3xSiTix [15,16] were not observed in the present study. In all our alloys annealed at 900  T ( C)  1000, binary or ternary, we find the orthorhombic modification as the stable modification of Ti5Si4. The lattice parameters determined for a binary alloy Ti55.6Si44.4 annealed at 950  C are a ¼ 0.65064(3) nm, b ¼ 1.26779(7) nm, and c ¼ 0.66409(4) nm, which is in good agreement with the literature data [27]. The XRD intensities observed are compatible with the space group Pnma and match reasonably well the intensities calculated by assuming a Sm5Ge4-type structure [32]. This phase is considered to be the low temperature modification of Ti5Si4. In Table 3 the label Ti5Si4(l ) is used. The tetragonal modification of Ti5Si4 occurred in as cast alloys and thus should be labeled Ti5Si4(h). The Zr5Si4type crystal structure for this phase is confirmed. Prolonged heat treatment at 900  C or 1000  C transformed Ti5Si4(h) into the orthorhombic modification Ti5Si4(l ). This reaction is very sluggish and might take forever at even lower temperatures. Thus, the metastable occurrence of the tetragonal modification Ti5Si4(h) at 850  C obviously caused the erroneous stability range assignment to this phase by Nickl and Sprenger [27].

37

1000  C

P63/mcm, Mn5Si3, D88

b

c

0.51459(3)

2.8 10.1 30.9

38.4 41.1 34.0

DTA signals ( C) on heating at a rate of 5 K/min Ti (at%) Si (at%) Fe (at%) Lattice parameters (nm)

Space group, prototype, Strukturbericht type

3.1. Crystal structures of the ternary phases

Ti (at%)

Phases present

3. Results and discussion

Si (at%)

Heat treatment

AG, Oberkochen, Germany; as cast alloys: model JSM-6400, JEOL, Tokyo, Japan). Thermal analysis was carried out in alumina crucibles under a stream of 99.999 wt% pure Ar between room temperature and 1500  C with a heating and cooling rate of 5 K/min using a DTA apparatus (model 701L, from Ba¨hr Thermoanalyse GmbH, Hu¨llhorst, Germany). The temperature was measured with PtePt/Rh thermocouples calibrated to the melting temperatures of Al (660.3  C), Ag (961.8  C) and Au (1064.2  C), Si (1413.8  C), and Ni (1455.2  C). The DTA signals read from the heating curves were used. Selected alloys were analysed up to 1700  C using a high temperature DTA from Setaram, Caluire, France (model SETSYS 18 DTA/ DSC).

3.2. Solid state phase equilibria

Fe (at%)

Table 3 (continued)

Trace Small U1 Offset

F. Weitzer et al. / Intermetallics 16 (2008) 273e282

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F. Weitzer et al. / Intermetallics 16 (2008) 273e282

measurements. In the isothermal section at 900  C (Fig. 2) equilibria among seven ternary phases were observed. These include the crystallographically already characterized phases t1-FeSi2Ti, t2-FeSiTi, and t3-FeSi3Ti, as well as the four phases t5, t7, t8, and t9, for which only the chemical compositions are determined so far (Table 1). In many alloys within the composition region t1et2eTi5Si3eTi5Si4 the large temperature gap between solidification and the temperature of the heat treatment caused the cores of the (relatively large) primary grains of Ti5Si3 metastably to remain and thus wrongly to pretend equilibration between this phase and t1. This made it extremely tedious to elucidate the true phase equilibria, because the ternary phases t7, t8, and t9 are close in composition and all three have very complicated XRD powder patterns without their crystal structures being known up to now. The situation was further complicated by the occasional occurrence of the (metastable) tetragonal modification of Ti5Si4(h) in addition to or instead of the orthorhombic equilibrium phase Ti5Si4(l ) due to insufficient equilibration. It is thus not just coincidental that the composition Fe10Si44Ti46 reported by Markiv et al. [7] for a phase X0 is the average of the compositions of t9 þ Ti5Si4(l ) and the composition Fe15Si40Ti45 reported for a phase X00 is the average of the compositions of t7 þ t8. Assuming each of these phase pairs to be just one phase each, the tie lines observed in the present study corroborate all tie lines reported for 800  C [7]. For the Si-rich phase t5 with the composition Fe7.4Si64.3Ti28.3 no corresponding observations were

279

made at 800  C, however. The phase Fe2Si5, which is stable in the FeeSi binary only above 937  C, apparently is stabilized at least down to 900  C by dissolving about 0.5e1 at% Ti. An additional ternary phase, labeled t4, occurs as a fourth phase in alloys having compositions within the phase triangle FeSi þ t1 þ t2. This phase still was found after anneal at 950  C but disappears upon annealing at 1000  C (Table 3). Since t4 was observed in as cast alloys (see below), it is considered to be a high temperature phase, which is stable only at T > 1000  C. At 1000  C the kinetics of its decomposition are apparently within the time scale of our annealing experiments, but t4 metastably remains even after prolonged annealing (four weeks) at 950  C or 900  C. The composition of t4 was determined as Fe28.1Si45.6Ti26.3 (Table 1). The Fe-rich phase Fe7Si2Ti [9] was not observed at 900  C in the present work, indicating the lower temperature limit of its stability to be above that temperature, too. All other tie lines of Fig. 2 match the tie lines at 1000  C as reported by Lo¨ffler [9]. For all of the ternary phases but t2-FeSiTi the homogeneity ranges measured by EDX do not exceed 1 at%. This corroborates the findings of Lo¨ffler and disagrees with Markiv et al., who observed a noticeable homogeneity range for t1, but none for t2. For the binary phase Fe2Ti extensive, and for Fe3Si, Fe5Si3, as well as Ti5Si3 limited solubilities of the third elements are observed (Fig. 2). Again, this agrees very well with Lo¨ffler. Because only few Si-poor alloys were

Fig. 3. Reaction scheme of the Si-poor region. Note: The temperature for reaction U10 is taken from the DTA data of alloy 713 in Table 4.9 of Ref. [9].

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investigated in the present work, the phase boundaries in the Si-poor region of Fig. 2 (i.e., for less than 25 at% Si) were determined by interpolation of the data reported by Lo¨ffler [9]. 3.3. Phase equilibria involving the liquid phase Gibbs phase rule requires for a ternary system that the formation of a fourth phase in a three-phase alloy is an invariant reaction. Since all-solid phase reactions would cause only a small DTA signal (as would do liquid forming reactions caused by a trace presence of residual non-equilibrium phases), the first strong peak due to an invariant reaction in our DTA heating runs of three-phase alloys was assigned to the formation of liquid phase from the three solid phases identified (Table 3). By coupling the resulting four phase equilibria with the corresponding onset temperatures of the

DTA signals the reaction schemes shown in Figs. 3 and 4 are obtained. These reaction schemes account for all DTA signals observed except one small peak at 1539  C in alloy Fe12.3Si43.9Ti43.8 and one peak at 1214  C in alloy Fe45Si15Ti40. In the Si-poor region (Fig. 3) most of the invariant reactions involving the liquid phase were recorded. Accepting the literature value [6] t2-FeSiTi is assumed to melt congruently above 1700  C, a temperature outside the range of our DTA instruments. The Laves-phase (Fe1xSix)2Ti saturated with Si is observed to melt incongruently at 1591  C, up from the (congruent) binary melting temperature of 1427  C [26]. This agrees with literature data [5] reporting that the substitution of Fe by Si increases the melting temperature of this phase. The peritectic formations of the ternary phases t4 (P5 at 1414  C) and t3 (P7 at 1241  C), as well as the ternary eutectics

Fig. 4. Reaction scheme of the Si-rich region.

F. Weitzer et al. / Intermetallics 16 (2008) 273e282

E1: L ¼ t2 þ t4 þ FeSi at 1254  C, E3: L ¼ t3 þ a2-Fe3Si þ Fe5Si3 at 1151  C and E4: L ¼ FeTi þ Ti5Si3 þ b(Ti) at 1034  C occur in this part of the reaction scheme. For comparing the present results with the invariant reactions reported in the literature it must be taken into account that Vogel and Schlu¨ter [3] accepted a different FeeSi binary diagram (without the phase Fe2Si) and that the phase t3 apparently did not form under the relatively high cooling rates employed (90 K/min). As a consequence, the reactions P7, P8, and U16 do not show up in their reaction scheme, and the reactions U12 (at 1199  C), U13 (at 1194  C), and E3 (at 1151  C) were described as L þ FeSi þ t2 ¼ Fe5Si3 (Fe3Si2) (P at 1200  C), L þ FeSi ¼ a2-Fe3Si þ Fe5Si3 (Fe3Si2) (U at 1180  C), and L ¼ t2 þ a2-Fe3Si þ Fe5Si3 (Fe3Si2) (E at 1145  C), respectively. For the latter reaction Lo¨ffler [9] observed a temperature of 1149  C (alloy Fe58.1Si35.3Ti6.6, annealed at 1000  C). The reaction U10 (at 1202  2  C; alloy Fe58Si28Ti14 of Ref. [9]) matches the transition reaction observed by Vogel and Schlu¨ter at 1185  C. Though melting temperatures along the monovariant eutectic trough between a,a2(Fe) and the Laves-phase (Fe1xSix)2Ti are observed to descend with increasing Si-content from 1317  C (alloy Fe75.4Si6.4Ti18.2 containing 6.4 at% Si; [9]) to 1311  C (alloy Fe69.1Si12.7Ti18.2 containing 12.7 at% Si; [9]), and 1302  C (alloy Fe70Si20Ti10; this work), these data confirm the reported temperature maximum at about 5 at% Si [3], because all these temperatures are above 1293  C, the eutectic temperature in the FeeTi binary [26]. Thus, the present reaction scheme corroborates very well all invariant reactions reported by Vogel and Schlu¨ter, except the peculiar, very Fe-rich ternary eutectic at 1205  C. The phase equilibria at 1150  C derived from the present reaction scheme match the isothermal section at 1150  C presented by Lo¨ffler [9] with the exception that we find a2-Fe3Si to coexist with Fe5Si3, while this author reported Fe2Si þ t3 in equilibrium. The latter tie line is based on EPMA data of a single alloy (alloy Fe58.1Si35.3Ti6.6). The same alloy equilibrated at 1000  C showed t3 þ a2-Fe3Si þ Fe5Si3. It cannot be ruled out that an all-solid state reaction a2-Fe3Si þ Fe5Si3 ¼ Fe2Si þ t3 occurs at 1000 < T ( C) < 1150  C, but at present we have no direct evidence for that. Thermodynamic modeling may prove as a decisive tool to clear up this issue. The Si-rich part of the system (Fig. 4) is dominated by the congruently melting phase t1-FeSi2Ti. The melting point of this phase is above 1500  C. It was not measured, but observed to be outside the range of our standard DTA instrument. The liquidus surface is characterized by a eutectic valley descending from the pseudobinary maximum emax1: L ¼ t2 þ Ti5Si3 at 1662  C through peritectic formations of t8 (P1 at 1640  C), t9 (P2 at 1597  C), t6 (P4 at  1480  C), and t5 (P6 at 1263  C) to end in the ternary eutectic E2: L ¼ t1 þ (Si) þ Fe2Si5 at 1175  C. The composition of the phase t6 was determined to be close to Fe12.5Si49Ti38.5 in the as cast alloy Fe10Si42Ti48, where it is observed as primary phase. This composition is considered to be tentative only, because in the present study t6 was not observed in any equilibrated alloy. The temperature of the lower stability limit for t6 must be substantially higher than 1000  C,

281

because we were not successful in quenching t6 to room temperature. The incongruent melting of t6 occurs at P4 above 1450  C, the temperature of the transition reaction U5, but most likely below 1480  C, the temperature of the transition temperature U4, although no DTA peak was observed between these temperatures in alloy Fe5Si47.5Ti47.5. The absence of such a DTA peak is interpreted to be due to slow nucleation of the new solid phase to be formed upon the incongruent melting of t6 in that alloy. Such kinetic hindrance occurs frequently upon heating alloys having compositions between the liquid phase and the phase to melt incongruently [33]. Combining these reaction schemes with the observations of the primary crystallizing phase in as cast alloys, the liquidus projection shown in Fig. 5 is obtained. As can be seen, of the ternary phases only t1 and t2 have large regions of primary crystallization, while limited such regions occur for the phases t3, t4, t5, t6, t8, and t9. For the phase t7 no coexistence with the liquid phase occurs. The entirely ternary nature of the primary region of crystallization for Fe5Si3 [3] is confirmed. These results agree with the observation of the phases t1et6, t8 and t9 in as cast alloys (not necessarily as primary phase, however), while the phase t7 could not be observed in such alloys. As in the present work the iron-rich region (labeled a,a2(Fe)) was not investigated with great detail, no information about the coexistence of Fe7Si2Ti with the liquid was collected. 4. Conclusions The constitution of the system FeeSieTi was investigated over the entire composition range between 900  C and 1700  C. Nine ternary phases are observed. The crystal structures for the phases t1-FeSi2Ti and t2-FeSiTi are corroborated. The crystal structure of the phase t3-Fe4Si3Ti is newly

Fig. 5. Liquidus projection for FeeSieTi. The phase fields of a(Fe) and a2Fe3Si are not separated and the term a,a2(Fe) is used. Due to space limitations U-type (transition) reactions are not labeled.

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determined to be isostructural with Pd40Sn31Y13. The phase equilibria at 900  C are established. In this isothermal section seven of the nine ternary phases (all but t4 and t6) are stable. The tie lines observed agree very well with previous literature. A reaction scheme linking the equilibria observed in equilibrated alloys (at 900  C or 1000  C, respectively) with the liquid phase is proposed. The resulting liquidus projection shows large areas of primary solidification for the phases t1 and t2, as well as small areas of primary crystallisation for the phases t3et6, t8, and t9. The phase t7 was not observed to form directly from the liquid phase.

[11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

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