Phase equilibria at 1000°C in the AlCSiTi quaternary system: An experimental approach

Phase equilibria at 1000°C in the AlCSiTi quaternary system: An experimental approach

MATERIALS SCIENCE & ENGINEERING E L S EV I E R Materials Science and EngineeringA229 (t997) 95-113 A Phase equilibria at 1000°C in the A I - C - S ...

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MATERIALS SCIENCE & ENGINEERING E L S EV I E R

Materials Science and EngineeringA229 (t997) 95-113

A

Phase equilibria at 1000°C in the A I - C - S i - T i quaternary system: an experimental approach J.C. Viala *, N. Peillon, F. Bosselet, J. Bouix Laboratoire des Multimat~riaux et Interfaces, UMR CN.RS No. 5615, Universit~ Lyon 1, 69622 Villeurbanne Cedex, France

Received 20 February 1996

Abstract

More than 50 powder mixtures with different compositions in the AI-C-Si-Ti quaternary system were cold-pressed, heated for 400 h at 1000°C under 105 Pa argon (1 atm) and rapidly cooled. Phases present in the as-treated samples were characterized by X-ray diffraction, optical metallography, scanning electron microscopy and electron probe microanalysis. From the results obtained, a three-dimensional isobaric-isothermal section of the AI-C-Si-Ti phase diagram has been constructed, using the classical equilateral tetrahedron representation. Different two-dimensional sections through this tetrahedron and projections on its faces have also been drawn. These constructions are discussed with the aim to provide a comprehensive description of the phase equilibria that tend to be established in the AI-C-Si-Ti quaternary system at 1000°C under a pressure of 1 atm. Among the most important results, one can mention the existence of liquid-solid phase equilibria between AI-Si base liquids and each of the solid phases TihSi3Cx, Ti3SiC2 and TiC. The Si contents of the A1-Si base conjugate liquids are, respectively, within the limits 8-10.5 at.%, 1-19 at.% and 0-11.5 at.%. Moreover, it is worth noting that TihSi3Cx conjugates with all the solid compounds of the system (A14C3 and SiC excepted) and that the three-dimensional isobaric-isothermal section is separated into two distinct parts by a three-phase equilibrium TisSi3Cx-A13Ti-TiC. The discussion also deals with the thermodynamic aspect of the high-temperature chemical interaction between siticon carbide SiC and Ti-A1 alloys. © 1997 Elsevier Science S.A. Keywords: Chemical reactions; Phase diagram; Phase equilibria; Silicon carbide; Titanium aluminides

1. Introduction

Detailed investigations have been conducted for several years on intermetallic matrix composites for applications in advanced aerospace systems such as new gas turbine engines or future hypersonic vehicles [1-3]. Despite these efforts, it has not yet been possible to overcome all the problems and to achieve all the requirements for reliable high-temperature structural materials. For most of the intermetallic matrix composite systems studied, it seems that progress has been slowed down by the lack of basic knowledge on the physical, mechanical and chemical behaviour of the matrix-reinforcement interface. Due to this lack, it has often been difficult to get the thorough understanding of the experimentally observed phenomena, which is needed for developing useful models. For example, as * Corresponding author. Tel.: + 44 72431897; fax: + 44 72440618. 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PTT ~ t ~ O 9 ] _ ~ C l O g ( Q T ~ l f ) 0 9 _ ~

concerns the chemical behaviour of the interface, it is clear that understanding the chemical interaction processes occurring at high temperatures during the elaboration or use of these composites implies a detailed knowledge of the phase diagrams of the relevant systems. However, most of these phase diagrams are only partially known, even unknown. In this context, the present work focused on the A 1 - C - S i - T i phase diagram has been undertaken to provide the grounds for a better understanding of the interface chemistry of a class of a priori very promising composite materials based on intermetaltics: composites made of a titanium aluminide matrix reinforced with silicon carbide fibres. As thermodynamic calculations were not possible because accurate data on the free energy of formation of several compounds were not available, it is by isothermal diffusion experiments that we have approached the phase equilibria in the A t - C - S i - T i quaternary system at 1000°C under t atm.

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J.C. Viafa et al./ Materials Science and Engineering A229 (1997) 95-113

2. Phase equilibria in the four ternary subsystems adjacent to the A I - C - S i - T i quaternary system Prior to the determination of the phase equilibria in the A1-C-Si-Ti quaternary system, it was necessary to collect reliable data on the phase equilibria existing at 1000°C under 1 atm in the four adjoining ternary subsystems AI-C-Si, C-Si-Ti, A I - C - T i and A I - S i Ti. This was achieved by reviewing literature data. At this early stage, however, it appeared that simplifications had to be made in order to maintain sufficient clarity of further representations of the A I - C - S i - T i quaternary phase diagram. Consequently, not very stable compounds playing a minor part in the ternary phase diagrams were omitted, compounds for which a narrow or undefined homogeneity range was reported were represented by a point and compounds for which a wide homogeneity range was reported were schematically represented by a line segment. 2.t. The A I - C - S i

C (axes in at-%)

system

High-temperature phase, equilibria in the A I - C - S i ternary system have been experimentally determined by Oden and McCune [4] and recalculated by Lukas [5]. According to these authors, two ternary compounds A18SiC7 [6] and A14SiC4 [7] are stable below about 2085 and 2075°C, respectively. These ternary compounds must then figure in the stable A1-C-Si ternary phase diagram at medium or low temperature. However, experiments have shown that formation of these ternary compounds by reaction between either the elements or the binary carbides SiC and A14C3 only occurs at temperatures higher than about t200 or 1400°C, respectively [8]. Hence, if these ternary compounds are not initially present at t000°C or below, metastable phase equilibria excluding them are to be observed in the A I - C - S i system [9]. It is these metastable phase equilibria which are represented in Fig. 1 and will be taken into account in t h e following. It is recalled that at 1000°C, a liquid phase extends along the A1-Si segment from 0 to 44 at.% Si and that the AI-Si liquid simultaneously in equilibrium with SiC and A14C3 contains 13 + 0.6 at.% silicon [9] and about 30 at. ppm carbon [10]. 2.2. The C - S i - T i

Ti0.48Sio.ztC0.31 [1t], Tio.49_0.51Si0.15_0.17C0.33_0.36 [12], Ti0.46Si0.asC0.35 [16] and Ti0.4sSi0.16C0.3~ [17]. More generally, all the phases reported in Fig. 2 have been considered as pure elements or as perfectly stoichiometric compounds, except for titanium carbide (referred to as TiC) and for the ternary solid solution deriving from the binary silicide TisSi3 (referred to as TisSi3Q~). Titanium carbide TiC has been represented as a line compound containing 35-49 at.% C, according to data reported by Murray [18]. The ternary solid solution TisSi3Cx has also been represented as a line compound: to comply with the experimental results we have obtained and with the data reported by Brukl [11] and

AI

L

Sl

Fig. 1. Simplifiedisothermalsectionat I000°C of the AI-C-Si pfiase diagram. C

I

axes in at'%

/~

system

The C - S i - T i section at 1000°C under 1 atm presented in Fig. 2 has been derived from experimental results obtained at 1100-1300°C [11-13]. To simplify the representation, the ternary carbide Ti3SiCz has been represented by a point having the composition of the perfect crystal [14,15], although the existence of a homogeneity range for this phase is very likely, as suggested by the different chemical' formulae reported:

13Ti

TisSiaC x TiSi

TiSi 2

Fig. 2. SimplifiedC-Si-Ti section at 1000°C,

Si

J.C. Viala et a l . / Materials Science and Engineering A229 (1997) 95-113

compound (P-phase) with the chemical formula Ti3A1C [19] and a hexagonal compound (H-phase) with the chemical formula Ti2A1C [20]. Fig. 3 shows the simplified A I - C - T i isotherm at 1000°C constructed from these data and from a review by Hayes [27]. Simplifications concern the binary phases A12Ti and A15Ti2, which have been omitted, and the homogeneity ranges for Ti2A1C and Ti3A1C, which have been neglected. In fact, the only homogeneity ranges represented at 1000°C in Fig. 3 correspond to the binary phases TiC (35-49 at.% C), TiA1 (41.5-52 at.% Ti), Ti3A1 (63-77 at.% Ti), e-Ti (80.5-89.5 at.% Ti) and /%Ti (92-100 at.% Ti): values between brackets were taken from Refs. [18,27,28]. Concerning the liquid phase located in the A1 corner, it will be recalled that it can contain up to 1.5 at.% titanium [28] and about 30 at. ppm carbon [10].

$1 axes in at-% "~

--

/

//~

~,

1\ /g, si2

.a.I

i

'AlaTi'

1

TiAi

I

97

i

TIaAI aYi'

fm

2.4. The A l - S i - T i system

Fig. 3. Simplified A I - C - T i section at I000°C.

O

AI4C3/£-- ' - ' - - ~

/

AI

'AlaTi'

~~TiC

//P'//I,

TiAI

Ti3AI

-TkAIc

aTi lrn

Fig. 4. Simplified A1-Si-Ti section at 1000°C.

Maline et al. [17], it has been assumed that the carbon content of this phase can vary from 0 to about 10 at.% while its silicon content remains constant. Finally, it is to be noted that the silicides Ti3Si and TisSi4 have been omitted in Fig. 2 since they are not stable in the presence of carbon and will, therefore, play a minor part in the A 1 - C - S i - T i phase equilibria.

The simplified A1-Si-Ti section at 1000°C under t atm reported in Fig. 4 is based on experimental results obtained by Schob et al. [29], by Raman and Schubert [30], and on a recent literature survey by Perrot [31]. Two ternary compounds were unambiguously characterized in this system: TiTA15Si12, stable only at temperatures lower than 900°C [30] and Ti(AlxSil-x)2, stable up to at least 1200°C [29]; the existence of a third ternary compound was suggested [29] but not further confirmed. Consequently, only the ternary compound Ti(AlxSit_ x)2 has been reported at 1000°C in Fig. 4. As the literature data concerning the composition at 1000°C of each of the phases involved in the three-phase equilibrium A13Ti-Ti(AlxSit_x)2-L were not precise enough, an experimental determination was carried out. From the analytical results reported in Table 1, it is clear that at 1000°C, 14% of the A1 atoms constituting the aluminide phase A13Ti can be replaced by Si atoms and that the ternary silicide Ti(AI~ Sit_ ~)2 can accommodate up to about 20 at.% A1 (i.e., x = 0.3). Accordingly, the phase designated as A13Ti is represented as a line compound with a constant Ti content of 75 at.% and an Si content varying from 0 to 14 at.%; as for the ternary silicide designated in the following as Ti(A1,Si)2, it is represented as a stoichiometric compound having the composition At:Si:Ti -- 20:47:33 (in at.%). Finally, it will be recalled that, consistently with the three other ternary sections drawn at 1000°C, the binary phases Ti3Si, TisSi4, A12Ti and A15Ti2 have been omitted.

2.3. The A I - C - T i system

Compound formation and phase equilibria in the A1-C-Ti ternary system have been experimentally studied by several authors [19-26]. Two ternary phases were characterized in this system: a perovskite-type

3. Representation of the phase equilibria in the A1-C-Si-Ti quaternary system The most common figure used to represent phase equilibria in a quaternary system at a fixed pressure P

J.C. Viala et at./Materials Science and Engineering A229 (1997) 95-I13

98

Table 1 Reaction a t 1000°C for 400 h in the AI-Si-Ti ternary subsystem: composition of the phases involved in the three-phase equilibrium A13Ti-Ti(A1,Si)2-L

Sample no.

I

Initialmixturecomposition (at.%) A1

Si

Ti

60

20

20

- A m / m o (%0) Phases characterized by XRD

25

Phases composition as determined by EPMA (at.%) A1

Si

Ti

88.5 60.6 18.8

10 14.2 48.6

~1.5 25.2 32.6

L = AI+Si A13Ti (blocks) Ti(AI,Si)2

Table 2 Interactions (t000°C, 400 h, 1 atm) in the A13Ti-SisC triangle: crystalline nature (XRD results) and composition (EPMA results) of phases formed in samples 2, 3 and 4 Sample no.

-~n/mo) (%) Phases characterized by XRD

Initial mixture composition (at.%)

0.75Al+0.25Ti

0.5SIC

Si

C

2

30

28

0

42

8

3

30

35

35

0

19

Ca

A1

Cgr. SiC TiC AI4C3

-----

---__

L = AI+Si Si

---

---

Ti(A1,Si)2 4

50

30

20

0

0.6

Phases composition as determined by EPMA (at.%)

1

[

SiC

--

--

L = AI+ Si SiC Ti(A1,Si)2 Ti3SiC2

ND 50 ND 33.8

79.5 ND 13.5 1.2

Si

Ti

19 50 54 17

~,I.5 ND 32.5 48

ND, researched but not detected; - - , not determined. Determined by difference with an estimated accuracy of + 3 at.%.

and at a given temperature T is an equilateral tetrahedron, the four apices of which correspond to the pure components [32]. In this figure, which is a three-dimensional isobaric-isothermal section through the phase diagram of the system, the triangular faces represent isobaric-isothermal sections of four ternary subsystems. Accordingly, to represent the phase equilibria at 1000°C under t atm in the A I - C - S i - T i quaternary system, we have constructed using wood and metallic rods an equilateral tetrahedron having A1, C, Si and Ti as apices and the four ternary sections at 1000°C drawn in Figs. 1-4 as faces: in the following, this tetrahedron, o f which a general view will be given in Fig. 17, is designated as the 1000°C A1-C-Si-Ti isotherm. Such a spatial model being difficult to represent in two dimensions and to exploit, four two-dimensional sections of special interest (A13Ti-Si-C , T i 3 A l - S i ~ C , TiAI-Si-C and AI-Ti-SiC) have been drawn through it to show the phase relations and the phase fields boundaries. The edges of the triangles and quadrilaterals figured in these sections were constructed by trian-

gulation. In the following, these two-dimensional sections will be designated as isopleths. Another mode to give a two-dimensional representation of quaternary phase equilibria consists in projecting characteristic lines of liquidus, solidus or solvus surfaces of the tetrahedron from an apex onto the opposite face. Such conic projections will also be used to represent the arrangement of the phase fields with each other. It is important to keep in mind that for condensed phase equilibria in a quaternary system under a constant pressure and at a given temperature, the phase rule is expressed by the equation v = C~-f, where v is the variance, C the number of components (C = 4) and f the mwnber of phases that coexist at equilibrium. Consequently, the following points hold: (1) A four-phase equilibrium will be invariant, which means that the composition of each of the four coexisting phases will be fixed. Such an invariant equilibrium will be represented in the 1000°C A I - C - S i - T i isotherm by a tetrahedron, the four apices of which will correspond to unique and precisely defined composi-

J.C. Viala et a l . / Materials Science and Engineer#~g A229 (I997) 9 5 - 1 t 3

99

Table 3 Interactions (1000°C, 400 h, 1 atm) in the A13Ti-Si-C triangle: crystalline nature (XRD results) and composition (EPMA results) of phases formed in samples 5-9 Sample no.

5 6 7 8 9

-Am/m o

Initial mixture composition (at.%)

Phases characterized by XRD

Phases composition as determined by EPMA (at.%)

(%)

0.75A1+ 0.25Ti

0.5SIC

Si

C

30 40 50 70 50

70 60 50 30 40

0 0 0 0 0

0 0 0 0 10

22 5 8 5" --

L = AI+Si Ti3SiC2 SiC AI4C3

C~

AI

Si

Ti

ND 33.6 50 43.6

86.5 1.6 ND 54.3

12 15.1 50 2.1

~ 1.5 49.7 ND ND

ND, researched but not detected; - - , not determined. Determined by difference with an estimated accuracy of + 3 at.%. Tabte 4 Interactions (t000°C, 400 h, 1 atm) in the AI3Ti-Si-C triangle: crystalline nature (XRD results) and composition EPMA results) of phases formed in samples 11-16 Sample no.

-Am/m o (%)

Initial mixture composition (at.%)

0~75A1+ 0.25Ti

0.5SIC

Si

C

11 12 13 14

60 74 74 72

12 19 0 0

0 0 9.5 8

28 7 t6.5 20

9 1.6 ---

15 16

75 73

0 0

5 4

20 23

---

Phases characterized by XRD

Phases composition as determined by EPMA (at.%) Ca

AI

Si

Ti

L = Al+Si Ti3SiCa TiC A14C3

ND 30.3 49.8 46.7

89.5 2.3 ND 48.6

~ 10 16.7 ND 4.7

,,~ 1.5 50,7 50,2 ND

L = AI+ Si TiC

---

---

---

A14C

--

__

__

3

ND, researched but not detected; - - , not determined. Determined by difference with an estimated accuracy of _+3 at.%.

tions of each of the four phases involved. In plane sections through the 1000°C A I - C - S i - T i isotherm (isopleths), this equilibrium will appear either as a triangle or a quadrilateral. (2) A three-phase equilibrium will be monovariant (v = 1): the three-phase volume representing this equilibrium will then be constituted by an array of tie triangles, the apices of which will follow monovariant lines of the liquidus, solidus or solvus surfaces limiting each of the three phases involved. In many cases, however, only the variation in composition of two or one of the three phases involved is sufficient to be represented. In these cases, the two triangles share a corner or an edge. When no significant variation can be observed in the composition of any of the three phases involved, the volume corresponding to the three-phase equilibrium is reduced to a simple triangle. In the isopleths drawn through the three-dimensional isotherm, the three-phase equilibria will appear as regions generally limited by two straight lines, these lines being superposed in the latter case. To distinguish more easily these three-phase regions, some sections of the tie triangles have been schematically represented.

(3) For two-phase equilibria, the variance is equal to 2. Such two-phase regions will appear as volumes limited by curved surfaces in the three-dimensional isotherm and areas limited by curved lines in the twodimensional isopleths.

4. Experimental procedure As said beforehand, the phase equilibria that tend to be established at 1000°C under 1 atm in the A I - C - S i Ti system have been approached by isothermal diffusion experiments. For that purpose, powder mixtures with compositions lying mainly in the two-dimensional sections A13Ti-Si-C and Ti3AI-Si-C were prepared from commercial powders of aluminum (Fluka, purity over 99 wti%, grain size d < 150 gm), titanium (Merck, purity over 98 wt.%, d < 150 gin), carbon (Le Carbone Lorraine, spectrographic grade graphite, d ~ 1 gin), silicon (Merck, purity over 99 wt.%, d < 50 gm), c~hexagonal silicon carbide SiC (Alpha Ventron, purity over 99.8 wt.%, d < 50 gm, mainly 6H and 8H polytypes) and cubic titanium carbide with a C:Ti atomic

J.C. Viala et aL / Materials Science and Engineering A229 (I997) 95-113

100

Table 5 Interactions (1000°C, 400 h, 1 atm) in the A13Ti-Si-C triangle: crystalline nature (XRD results) and composition (EPMA results) of phases formed in samples 17-26 Sample no.

0.75A1+ 0.25Ti

0.5SIC

17 18

82 84

10 6.4

19 20

78 79

10 10

21

74

22

23 24 25

26

--Am~too (%)

Initial mixture composition (at.%)

Si

C

Phases composition as determined by EPMA (at.%) Ca

A1

Si

Ti

8 9.6

-1.5

L = A 1 (no Si) Ti3SiC2 TiC A13Ti (blocks)

ND 31.0 51 ND

97.7 6.5 ND 70,8

~,0.8 12.3 ND 4.4

,~1.5 50.2 49 24.8

12 11

0 0

0.7 1.0

L = Al+Si Ti3SiC 2 A13Ti (blocks) TisSi3C x

ND 32.2 ND 9.9

90.5 1.5 63,2 7.8

8 t6.6 12,4 34.0

~ 1.5 49.7 24.4 48.3

10

16

0

0.9

L = AI+ Si Ti(A1,Si)2 Ti3SiC 2 TisSi3C ~

ND ND 36.7 10,3

88.5 14.5 ND 7.7

10.5 52.4 15.8 34.5

~ 1.5 33.1 47.5 47.5

78

10

0

12

1

L=AI+Si TiC Ti3SiC 2 AI3Ti (platelets)

ND 47.9 31.2 ND

96.5 ND 2,7 66.2

~2 ND 15,5 8.6

~.1,5 52.1 50,6 25,2

9 88

4 0

0 2

0 10

0.4 2.6

L = AI (no Si) A13Ti (blocks TiC

. .

. .

. .

. .

.

.

.

. ~2

~1.5

80

70

20

15

0 0

Phases characterized by XRD

0

15

0

1.5

0

L=AI+Si

ND

96.2

Ti3SiC 2 A13Ti (blocks)

29.8 ND

3,6 64.8

16,2 9.9

50.4 25.3

L = AI+Si Ti(A1,Si)2 Ti3SiC a

ND ND 33.6

85.5 13.1 1.0

13 54,2 16.8

~ 1.5 32.7 48.6

ND, researched but not detected; - - , not determined. a Determined by difference with an estimated accuracy of _ 3 at.%.

ratio of 0.98 (Alpha Ventron, purity over 99 wt.%, d < 20 gm). The mixtures were then ball-homogenized for 10 min in a tungsten carbide mortar, cold-pressed under 240 MPa into small parallelepipedic rods weighing about 2.5 g (30 x 4 x 6 mm) and placed on alumina boats into silica tubes. After having sealed the silica tubes under approximately 0.3 atm of purified argon, they were heated first for 24 h at 700 +_ 5°C and then for 400 h at 1000 _+5°C, using a conventional horizontal tube furnace supplied by a stability controller. At the end of the isothermal treatment at 1000°C, the silica tube was driven out of the furnace and rapidly cooled in air: under such conditions, the cooling time constant for the sample was in the order of 1 min. On the one hand, preheating for 24 h at 700°C was necessary to avoid the selectivevaporization of aluminum during the early stage of the interaction. On the other hand, a heating time of 400 h appeared generally tong enough for completion of the reactions without excessive deviation to the initial composition due to aluminum vaporization. Finally, although not very fast, the cooling rate

appeared sufficient to freeze equilibria established during isothermal annealing at 1000°C. Characterization of the sintered rods thus obtained was carried out by different techniques. In a first step, the relative weight loss ( - Am/rno) was systematically controlled: it varied from about 2-4% for the densest samples to more than 20% for the most porous samples. One face of the treated rods was then grossly polished and characterized by X-ray diffraction (XRD) using a Philips PW 3020 goniometer (filtered Cu K~ radiation). When possible, sections of these rods were diamond-polished to a finish of 1 ~tm and examined by optical metallography (OM) and scanning electron microscopy (SEM). Further characterization of the phases produced was performed by electron probe microanalysis (EPMA), using a Cameca Camebax apparatus equipped with a wavelength dispersive spectrometer and an energy dispersive analyzer. An accelerating voltage of 10 kV, a regulated beam current of about 9 nA and a counting time of 15 s were selected as standard operating parameters. For each phase ana-

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J.C. I"iala et a l . / Materials Science and Enghzeerhzg A229 (1997) 95-113

lyzed, the counting rates simultaneously recorded for A1, Si and Ti in at least 10 different points were averaged, subtracted for background and referred to the counting rates recorded under the same conditions for pure A1, Si and Ti standards (also freshly diamondpolished to a finish of 1 #m). The ratios thus obtained were then corrected for atomic number, absorption and fluorescence using a classical correction program (ZAF), which led to AI, Si and Ti atomic contents with an estimated accuracy within _+ 1 at.%. After having verified that no foreign impurities were detectable in the fluorescence spectra of the analyzed sample (wavelength dispersive spectrometry for boron, nitrogen and oxygen and energy dispersive spectrometry for other elements with Z > 8), the carbon content was determined by difference with an estimated accuracy of + 3 at.%. Analyses of the AI-Si alloys resulting from the solidification of the liquid phase were performed either by EPMA or by XRD: for hypo-eutectic or near-eutectic alloys (less than 15 at.% Si), the composition was determined by EPMA using a particular procedure [9]; for hyper-eutectic alloys, the silicon content was estimated by comparing the intensities of the X-ray diffraction lines characteristic for A1 and Si in the samples and in cast AI-Si standards. After having determined the nature and composition of each of the coexisting phases in the samples equilibrated at 1000°C, the corresponding phase fields were constructed by degrees in the 1000°C A I - C - S i - T i isotherm. In the forthcoming presentation of the results, special attention will be paid to the A13Ti-Si-C and Ti3AI-Si-C sections through this isotherm: it is effectively in these isopleths that relations existing between the initial composition of a sample, the nature and composition of the phases present in it after reaction and the corresponding phase field appear most clearly.

tions were found to consist, after reaction, of the same four phases exhibiting the same characteristic X-ray diffraction peaks and having the same composition to the precision of EPMA, it was concluded that conditions for the establishment of an invariant four-phase equilibrium at constant pressure had been realized at 1000°C. Results concerning these samples were then grouped under the same line and, as the compositions of the phases were not significantly different from one sample to another, the contents reported are average values characteristic for the corresponding equilibrium. C

1 AI3T i 4

Sl

Fig. 5. Position in the A13Ti-Si-C triangle of the powder mixtures listed in Tables 2-5.

C

/ 5. Results and interpretation

~

~

1000°C

-

~

~ s [ o ,

8[0,TiC,AI4C3,T1

A14C'3T1, , L(11,,ff-12)~ ~ S : C ' 1 1 ' L ( I 2

TiC'AI403'TI'L(I'.5)~ ~;'i;:

AI4e3,T1,L~12) "19)

;::L(2'19'~;)

5.1. Phase equilibria at lO00°C in the AI3Ti-Si-C isopleth In a first series of experiments, mixtures with compositions located in the A13Ti-Si-C triangle, i.e., mixtures with an atomic ratio AI:Ti constant and equal to 3, were heated for 400 h at 1000°C and characterized. The initial atomic composition of these mixtures is given in the left part of Tables 2-5. Fig. 5 indicates the position of the corresponding point in the A13Ti-Si-C triangle. The relative weight loss ( - Arn/mo) measured for each mixture after heat treatment as well as the nature and the composition in at.% of each phase present in this mixture are reported in the right part of Tables 2-5. When several samples with different initial composi-

AI3Ti . . . . . . . . .

/; "r] T L

o

/ ~

L(O.t;"~/~ ~

TiC,AI:3"I"i,

$1

L(8)

Ti5SI3Cx,AI3TI,T1, v2, Loom~) ~ \ ' ~ / Z /--'ri5siaCx'TI'T2'L(IO'`) ~ ,-----Ti5SlaCx,T1, y//#'~/--TisSi3OX, AI3Ti,t-(8-10) //~Tis$I3Cx N3TI T2 ),

//--%,

L(8.I&5) L(lO)

AI3TI

Fig. 6. The 1000°C A13Ti-Si-C isopleth: (upper part) general view; (lower part) detailed view of the A13Ti comer.

102

J,C. Viata et at./Materials Science and Engineering A229 (1997) 95-I13

Fig. 7. Optical micrograph of sample 4 showing the phases involved in the four-phase equilibrium SiC-Ti3SiC2-Ti(A1,Si)2-L(19).

Results obtained for the first two mixtures figuring in Table 2 were easy to interpret, although only XRD characterization could be achieved (these samples were not dense enough for being polished and then observed by optical metallography or analyzed by EPMA). For sample 2, all the X-ray reflections observed were characteristic, in angular position and intensity, for the well-known crystalline phases C (graphite), SiC (c~hexagonal, 6H polytype), AI~C3 (hexagonal) and TiC (cubic, carbon-rich). As no extra reflection remained unindexed, it was concluded that an invariant fourphase equilibrium between these four phases had been reached at 1000°C. On the A13Ti-Si-C isopleth presented in Fig. 6, the section of the corresponding tetrahedron is the triangle with the C apex. For sample 3, X-ray diffraction revealed the presence of two elements A1 and Si, and that of two compounds, unreacted SiC in little amounts and Ti(A1,Si)2 in large amounts. As in the preceding case, no extra line remained unindexed. Taking into account the intensities of the diffraction lines characteristic for tile two elements, it was remarked that after cooling and solidification, sample 3 contained much more elemental silicon than elemental aluminum. As the solubility of silicon in liquid aluminum does not exceed 44 at.% at 1000°C, it was deduced that at that temperature, solid elemental silicon was in equilibrium with an A1-Si-based liquid containing 44 at.% Si, less than 2 at.% Ti and C traces

(according to literature data given in Section 2). Taking into account the Ti content of this liquid, the large amounts of Ti(A1,Si)z phase characterized by XRD in sample 3 could not have crystallized on cooling. It was then concluded that four phases were in equilibrium at 1000°C in sample 3 at the end of the heat treatment at 1000°C: SiC, Si, Ti(A1,Si)2 and an Al-Si liquid containing 44 at.% Si. In the A13Ti-Si-C isopleth, the selection of the tetrahedron representing this four-phase equilibrium is a triangle with Si as an apex. In contrast to mixtures 2 and 3, sample 4 could be polished and characterized by OM and EPMA. It can be seen in Fig. 7 that this sample consists of an AI-Si alloy matrix (AI-Si eutectic + Si block crystals) with an Si content of about 19 at.%, according to XRD results. Crystals of three sorts are embedded in this matrix: large faceted Ti(A1,Si)2 crystals with hexagonal shapes, dark unreacted SiC particles with jagged outlines and large round-shaped Ti3SiCz crystals often grown onto the SiC particles. Besides these major phases, small A13Ti platelets are also visible in the alloy matrix. Such platelets were in fact found in every sample in which a liquid phase was present. Very likely, these platelets precipitated on cooling from the liquid saturated ill titanium at 1000°C: analyses of the solidified alloys (A1 + Si + A13Ti platelets) revealed a titanium content of about 1.5 at.%, a value in agreement with the literature data summarized in Section 2. On the basis of these observations, we have concluded that four phases

Fig. 8. Optical micrograph of sample 8 showing the phases involved

in the four-phaseequilibriumSiC-A14C3-Ti3SiCa-L(I2).

J.C. Viala et at./Materials Science and Engineering A229 (I997) 9 5 - t i 3

103

based liquid containing 19 at.% Si which will be designated as L(19~. On the A13Ti-Si-C isopleth drawn in Fig. 6, the corresponding four-phase field is represented by a triangle with SiC as an apex. It is worth noting that the two four-phase tie tetrahedrons constructed from the results of experiments 3 and 4 both involve the same three phases: SiC, Ti(A1,Si)2 and an A1-Si liquid. In such a case, the law of adjoining phase regions [32] implies the existence, between the two tie tetrahedrons, of a three-phase region in which the three above-mentioned phases coexist. In the A13Ti-Si-C isopleth (Fig. 6), this three-phase region appears as a triangle having SiC as an apex. It is recalled that the thin lines drawn in this triangle represent sections by the A13Ti-Si-C plane of particular tie triangles in which SiC and Ti(A1,Si)2 are simultaneously in equilibrium with a liquid having an Si content ranging from 19 to 44 at.%, this liquid being designated as L(19-44))"

Fig. 9. Optical micrograph of sample 13 showing the phases involved in the four-phase equilibrium TiC-A14C3-Ti3SiCa-L(H.5).

L: r

Characterization by XRD and EPMA of samples 5, 6, 7, 8 and 9 after reaction at 1000°C for 400 h showed that they all contained the same four phases, although the starting mixtures had initially very different atomic compositions (Table 3). These phases are SiC, Ti3SiC2, AI4C3 and an A1-Si alloy containing 12 at.% Si. It can be seen in Fig. 8 that the remaining SiC particles with jagged outlines are either in direct contact with the

"

.

.

.

.

.

.

"

- 7 -

Fig. 10. Optical micrograph of sample 18 showing the phases involved in the four-phase equilibrium TiC-A13Ti-Ti3SiC2-Lm.

were in equilibrium at 1000°C in sample 4: three solid compounds SiC, Ti3SiC2 and Ti(A1,Si)2 and an A1-

Fig. 11. Optical micrograph of sampte 21 showing the phases involved in the four-phase equilibrium TisSi3C,-Ti3SiC2-Ti(A1,Si)2L(lo.5). In addition to the four major constituents, small dark needles are visible•

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d.C. ViaIa et al./ Materials Science and Engineering A229 (1997) 95-I13

TabIe 6 Phases characterized after heating TiC:AI-Si mixtures for I50 h at 1000°C (samples 27-33) Sample no.

27 28 29 30 31 32 33

Initial mixture composition (at.%) 0.5TIC

A1

Si

5 5 5 5 5 5 Ti = 2.5 C=2.5

95 89.3 85.5 84.55 83.6 81.7 86.45

0 5.7 9.5 10.45 11.4 13.4 8.55

Si:AI+ Si (at.%)

Phases characterized by OM, XRD and EPMA

0 6 I0 11 12 14 9

TiC, A1, (A1203, AI3Ti) TiC, Al-Si, (A1203, AI3Ti) TiC, Al-Si, (A1203, A13Ti) TiC, AI-Si, (AlaOs, A13Ti) A14C3, Ti3SiC2, AI-Si, (AI2Q, AIsTi) A14C3, TisSiC2, AI-Si, (AI203, A13Ti) TiC, AI-Si, (A14C3, AlzO3, AtsTi)

liquid phase or surrounded by numerous well-faceted TisSiC2 crystals and some small A14C3 crystals. On the basis of these results, a tie tetrahedron having SiC, A14C3, TisSiCa and an Al-based liquid containing 12 at.% Si as apices has been constructed in the A I - C - S i Ti isotherm at 1000°C. In the AIsTi-Si-C isopleth (Fig. 6), this four-phase region is represented by a triangle having SiC as an apex. It can be verified in Fig. 5 that the points corresponding to the composition of mixtures 5, 6, 7, 8 and 9 are all located inside the four-phase triangle SiC-A14C3-TisSiCa-L(I2~. Furthermore, application of the law of adjoining phase regions to the tie tetrahedrons SiC-TisSiCz-Ti(A1,Si)z-L(~9) and SiC-A14C3-TisSiC2-L(I:) implies the existence, between them, of a three-phase field in which SiC coexists with Ti3SiC2 and an Al-based liquid containing 12-19 at.% silicon. In Fig. 6, this three-phase field is referred to as SiC-TI-L(~z_19~. Another group of mixtures having different initial compositions but containing the same four phases with the same composition after 400 h reaction at 1000°C is composed of samples 11, 12, 13 and 14 (Table 4). In that case, the four phases are a carbon-rich titanium carbide that will be designated TiC, A14C3, TisSiCa and a near-eutectic AI-Si alloy. Determination by EPMA of the composition of the Al-based alloy present in these samples was not, however, possible, due to the fine dispersion of the other constituents (Fig. 9). It will be demonstrated in the following that the silicon content of the alloy simultaneously in equilibrium at 1000°C with TiC, A14C3 and Ti3SiCa is in fact of 115 _+ 0.5 at.%. According to these results, a tie tetrahedron representing the four-phase region TiCTi3SiCz-A14Cs-L(~Ls~ has been constructed in the A I - C - S i - T i isotherm at 1000°C. In the A13Ti-Si-C isopleth (Fig. 6), the section of this tetrahedron appears as a quadrilateral. At this stage of the construction, the existence of three other phase fields naturally derives from the application of the law of adjoining regions. These phase fields are the four-phase region SiC-TiCA14Cs-TisSiC2 and the two three-phase regions AI4C3-

Ti3SiC2-L(lt.5_12 ~ and TiC-A14C3-L(0_ll.s~, the latter phase field including points 15 and 16 (Table 4, Fig. 5). Results obtained for samples 17 and t8 (Table 5) show that they are both located in the same tie tetrahedron in the A I - C - S i - T i isotherm. Phases involved in this four-phase equilibrium are the three solid compounds TiC, A13Ti and TisSiC2 and quasi-pure liquid aluminum: in the A13Ti-Si-C isopleth (Fig. 6), this equilibrium corresponds to the triangle TiC-A13TiTi3SiCz-L(I~. The four constituents can be observed on the micrograph presented in Fig. 10. The metallic matrix is in this case an N-based alloy containing about 1.5 at.% Ti and about 1 at.% Si. The compounds A13Ti and Ti3SiC2 embedded in this matrix form large crystals (globular for A13Ti, well-facetted for TisSiC2). According to EPMA analyses, 4.4% of the A1 atoms of the A13Ti phase are replaced by Si atoms, while 6.5% of the Si atoms of the TisSiC2 phase are replaced by A1 atoms. As regards the phase designated as TiC, it has a composition very close to the 1:1 stoichiometry and appears under the form of small particles (1-3 ~tm in diameter) that are often concentrated with metallic aluminum in the core of TisSiC2 crystals. This particular morphology tends to indicate that a layer of TiC has formed around the SiC grains before their decomposition by liquid aluminum and that this TiC layer has served as nucleus for the growth of TisSiC2. Three other tie tetrahedrons have finally been constructed from the results obtained for samples 19, 20 and 21 (Table 5). In the A13Ti-Si-C isopleth (Fig. 6), the sections of these tetrahedrons correspond to the four-phase triangles TisSisCx-AlsTi-TisSiC2- L(8~ which includes points 19 and 20 (Fig. 5), TisSisCxTi3SiC2-Ti(A1,Si)z-L(10.5~ which includes point 21 (Fig. 5), and TisSisCx-A13Ti-Ti(AI,Si)2-L(t0} which naturally derives by construction from the two former triangles. The silicon content of the Al-based liquid involved in the latter four-phase equilibrium has been arbitrarily fixed at 10 at.%, which corresponds to the Si content of the liquid for the three-phase equilibrium A13TiTi(A1,Si)2-L(10~ in the AI-Si-Ti ternary subsystem at

105

J.C. Viala et al./Materials Science and Engineering A229 (1997) 95-113

1000°C (Table 1, Fig. 4). As no tie line exists in this subsystem between the silicide TisSi3 and the Al-based

C

TI3AI

Si

Fig. 14. Position in the Ti3AI-Si-C triangle of the powder mixtures listed in Table 7.

Fig. 12. Optical mierograph of sample 29 showing TiC particles in equilibrium at 1000°C with an AI-Si alloy containing 10 at.% Si.

liquid alloy, particular care was taken to characterize TisSi3Cx in samples 19, 20 and 21. In fact, the presence of TisSi3C x in these mixtures could be established with certainty by XRD. Moreover, OM and SEM observations (Fig. 11) showed that this phase was often in direct contact with the metallic matrix while EPMA analyses revealed that it contained about 10 at.% carbon and about 8 at.% aluminum. In addition to the four above-cited phases, small needles were also found in samples 19, 20 and 21 (and in these samples only). Due to their very small size, these needles could not be precisely characterized (some indications were obtained for A14C3). It is thought that these needles were not formed during the isothermal treatment but on cooling of the samples, as a consequence of a complex sequence of invariant transformations. Once these three last four-phase regions were established, it suffices to apply the law of adjoining regions for indexing the remaining three-phase regions and finishing the construction of the A13Ti-Si-C isopleth at 1000°C. Points 22, 23, 24, 25 and 26 were thus naturally included in the phase fields corresponding to their analyses by X R D and EPMA (Table 5, Figs. 5 and 6). 5.2. Phase equilibria at I000°C between TiC and A l - Si alloys

Fig. 13. Optical micrograph of sample 31 in which TiC particles were decomposed into A14C3 and Ti3SiC2 by reaction at 1000°C with an Al-Si alloy containing 12 at.% Si.

Analyses performed on samples 11-18 and 22 showed carbon-rich TiC to be in equilibrium at 1000°C not only with quasi-pure aluminum but also with A1-Si alloys. In order to get more detailed information about the composition of these alloys, samples containing small amounts (5 at.%) of C-rich TiC powder (Alpha

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J.C. Viala et at./Materials Science and Engineering A229 (1997) 95-1t3

Table 7 Interactions (1000°C, 400 h, 1 atm) in the Ti3AI-Si-C isopleth: crystalline nature (XRD results) of phases formed in samples 34-51 Sample no.

34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51

Initial mixture composition (at.%)

-Am/m o

Phases characterized by XRD

0 0 0 0 0 0 2.1 0 0.3 1.0 0.3 1.3 0 0.4 4.0 --1.2

TisSi3C~, Ti3AIC, Ti3A1 Ti2A1C, TisSi3Cx, Ti3AIC, Ti3A1, (TIC)

(%)

0.25Al+0.75Ti

0.5SIC

Si

C

84 80 78 78 72 72 65 65 65 61 73.5 69.5 65.5 63 50 38.5 23 40

16 20 22 0 0 0 0 0 35~ 0 0 0 0 0 0 24~ 24~ 60

0 0 0 6.2 7 4 14 4 0 12 17.5 21.5 25.5 33 44 37.5 53 0

0 0 0 15.8 21 24 21 31 0 27 9 9 9 4 6 0 0 0

TiC, Ti2A1C, TisSi3Cx, Ti3A1C TiC, A13Ti, Ti2AIC, TisSi3Cx, (C in samples 40 and 41)

TiC, AI3Ti, Ti3SiC2, TisSi3Cx TisSi3C~, Ti3A1, TiA1, Ti2A1C TisSi3Cx, TiAI, Ti2A1C TisSi3C~, TiAI, Ti2A1C, AI3Ti TisSi3Cx, Ti(A1,Si)a, AI:Ti (out of equilibrium) TisSi3C.~, Ti3SiC2, Ti(AI,Si)2 SiC, Ti(AI,Si)z, Ti3SiC2 SiC, Si, Ti(AI,Si)z; TiSi2 SiC, Ti3SiC2, Ti(A1,Si)2, L = Al+Si

- - , Not determined. a SiC ball-milled to a grain size of approximately 1 ~tm.

Ventron, purity over 99 wt.%, d < 20 gm) and a large excess (95 at.%) of A1 + Si powders were cold-pressed and annealed for 150 h at 1000°C under 1 atm argon. After rapid cooling, these samples were characterized by XRD, OM and EPMA: Table 6 summarizes the resets obtained. On the one hand, the major constituents characterized after heat treatment in samples 27-30, i.e., samples with an Si:A1 + Si atomic ratio lower than or equal to 11 at.%, were the C-rich TiC particles initially introduced and an AI-Si alloy with the same composition as that of the starting mixture. Fig. 12 shows the typical morphology of such samples. On the other hand, it appeared that the TiC particles had disappeared and were replaced by AlgC3 and Ti~SiC2 crystals in samples 31 and 32, i.e., samples with an Si:A1 + Si atomic ratio greater than or equal to 12 at.% (Fig. 13). In that case, the composition of~ the AI-Si alloy in large excess had not varied significantly. Finally, TiC was observed to form from the elements Ti and C in the presence of an AI-Si alloy containing 9 at.% Si (sample 33). It is clear from these results that C-rich TiC is in equilibrium at 1000°C with AI-Si alloys containing up to 11.5 at.% Si and that the silicon content of the AI-Si alloy simultaneously in equilibrium at 1000°C with TiC, AlgC3 and Ti3SiC2 (fourphase equilibrium) is 1t.5 _+0.5 at.%. A1203 crystals and small A13Ti platelets were also characterized as minor constituents, in all the samples of this series. The presence of A1203 may be attributed to an oxygen contamination of the starting A1 and TiC

powders, while that of small A13Ti platelets indicates that the liquid alloy equilibrated at 1000°C contains little amounts of titanium. As A13Ti platelets were found in higher density in samples 31 or 32 (Si-rich) than in samples 27 or 28 (Si-poor), it can be assumed that the maximum solubility at t000°C of Ti in Al-Si alloys, previously estimated to be about 1.5 at.%, increases with the Si content of these alloys.

5.3. _Phase equilibria at 1000°C in the Ti~Al-Si-C isopleth A third series of heat treatments was realized at 1000°C for 400 h under 1 atm a(gon on powder mixtures having a composition located in the Ti3AI-Si-C triangle, i.e., mixtures with an atomic ratio Ti:A1 constant and equal to 3. Fig. 14 shows the positions of these mixtures in the Ti3A1-Si-C triangle. The relative weight loss and the nature of the phases that were characterized by XRD after reaction are reported in Table 7. As only solid-state reactions occurred in this third series of experiments, characterization by EPMA was not possible, as the crystals formed were too small in size. Exploitation of the results reported in Table 7 was carried out using the same rationale as that outlined in Section 5.1. This rationale will not therefore be detailed here. It will, however, be remarked that some of the solid-state reactions occurring in this third series of heat treatments did not go to completion. In that case, the

107

J.C. Viala et aL / Materials Science and Engh~eering A229 (1997) 95-II3

incompletely reacted phase appears between brackets in Table 7 (it can be seen that the out-of-equilibrium phases are often C or TIC). To overcome this problem, some mixtures were prepared from an SiC powder with a very small grain size (d ~-, 1 gin). Another difficulty was the characterization by XRD of the silicides TiSi, TiSia and Ti(A1,Si)z in mixtures where they were present in low atomic fractions (weak characteristic lines and overlapping problems). Despite these difficulties, results reported in Table 7 appeared sufficiently self-consistent to permit the unambiguous construction of the lower part (C-poor compositions) of the Ti3AI-Si-C isopleth presented in FigL 15. As for the upper part of this isopleth (C-rich compositions), it was established by simple construction from the results previously obtained in the A13Ti-Si-C triangle. One can effectively fred in this upper part almost all the phase equilibria represented in the A13Ti-Si-C isopteth: the phase boundaries are distotted since the section plane has changed, but the arrangement of the phase fields is strictly preserved.

C

~;:

1000°C

;[140c;T::.L;lI'5)"~\~IJ~'~"'--SIC';:'L;:lo~51;)

TIAI . . . . . . . . . Ct /~ ,

TiC, AI3Ti, T1, L(1)

2[)', ~4",, TIC, AI3T1,T1, TiSSi3Ox- - x ~--;~ " ~ ' ~ & " ,

~._'~

TIC,N3TI,L(0.1)

TiC,AlaTi,"n2A~O~

~

Am, AI3Ti, TI2A~O ~

SI / / / / / /

///, / /////

N3Ti' TI' L(I"8) "~%C×,Tv L(e-lO.S) TIsSI3Cx,TI, T2, L(IO5) TtSSi3Cx,AI3TI, Zl, L i ; Ti5Si3Cx' Aj3~' Zl

~ .,~/..//~__TisS~Ox

"H/ ~ "~'(-
AITt,AI3Tt.TI5SlaCX- - - - @

1

aI~Ti,T~Si

~ "~\\\\/'~/~TiS$i3CX'NaTI'T2'L(Io)

/.J---'~F ~ ~ ' Y - " - - - - ~ % \ ' . ~, ~

"nAI

,; sl

Fig. 16. The 1000°C TiAI-Si~C isopleth: (upper part) general view; (lower part) detailed view of theTiA1 comer.

lOOOOC' TIC, At403, TI, L(I 1.5) IC, AI403, TI, L(12) SIC, T1, L(t2.19)

/ ~

. . ~ k ~ . . . . . - F - - T[SSlaOx, T1, T2, L(10.5) ~

TisS[3Cx, T 1, T2, T]Si2 _.~--TIsSi3Cx, T2, TtSI,AI3T1 ~_~._-.~, TIsSI3Cx,TISI, T2, TISi2

1

Sl

"riO, AI4O8, T 1, L(11.5) TiC, AI403, L(O.11.5)

~c, ,%T~,L(O-I) TIC, AI3Tf,TI2AIC, TisSi3Ox- - ~

/

~/~ " ~

L(1.11.5) ~O, AI3TI, T 1, L(1) TiC, T1,

~--~'~ ~

\'~,

/-~

TiC, AIsT1, T1, TISS[3Cx AJ3"n.TvL(1-8)

10.5)

TIC, T12AIC,Ti58t3Cx TiC, TI2AIO,TIsSI3Cx,T I 3 N C ~ Ti2AIC, ~58130×, Tt3AIC, T I 3 ~ . ~ _ ~

\'~,,/x'~ ~ssl3°,, T1, ~J3~,L8 \ \\\\~./,j--~ssl3c×, v~,~£n

\k~,Tissi~Ox,

\~-

Ti58130x, TISAI

"n3A,,"ro~,,Ti2A,c

~s%%*13~,L(o-10)

~ T i g S f 3 C x , AI3Ti .~" ~ ~ " TI58,3Cx, AI3TI,TiSt

¼,.rl3AI

Fig. 17, Photograph of the three-dimensional 1000°C A 1 - C - S i - T i isotherm (1 atm), as constructed from the experimental results obtained in the present work (real edge length: 400 mm).

_.5 "I]3AI,T~I,- = TIAI,

Tisst3ox

i \_..2

~ssloc×

TIN, "~5Si3Cx,

~a-n

~

" SI

Fig. t5. The 1000°C Ti3Al-Si-C isopleth: (upper part) general view; (lower part) detailed view of the Ti3A1 comer.

5.4. The IO00°C T i A I - S i - C isopleth

A two-dimensional section of the 100°C A 1 - C - S i - T i isotherm by the TiAI-Si-C plane is presented in Fig. 16. On this third isopleth, which naturally derives from the foregoing constructions, one can find the same phase field arrangements as those previously encountered in the A13Ti-Si-C and Ti3AI-Si-C isopleths (Figs. 6 and 15, respectively).

108

J.C. Viala et at./Materials Science and Engineering A229 (1997) 95-113

6. Discussion

O

In the following discussion, we will try, firstly, to provide a general description of the condensed phase equilibria that tend to be established in the A1-C-SiTi quaternary system at 1000°C under a pressure of 1 atm. Secondly, we will point out the thermodynamic principles governing the chemical interaction at 1000°C between silicon carbide SiC and AI-Ti alloys. 6.1. Phase equilibria at IO00°C in the A I - C - S i - T i system

A photograph of the three-dimensional representation of the 1000°C A1-C-Si-Ti isotherm, as it has

C

// ,

//

i

.98

I ~ , q ?

AI4C"~

~' $i

.

.

.

.

.

.

corner.

C A ~

.

.

.

.

.

Fig. 18. Partial view in true perspective of the three-dimensional 1000°c A I - C - S i - T i isotherm showing the solid-liquid phase equilibria.

Ti3SiC2 "~Ti

Fig. 20. Partial view of the three-dimensional 1000°C A I - C - S i - T i isotherm showing the solid-solid phase equilibria in the C-rich

~]

.

../

C~ ~ ',

T2=Ti(AI'Si)2 T1=TiaSiC2

S~/~"TiSi2",

issi3Cx',

AI ~ ~ " / / ~ ' - / / ~

"'''-,-\',

~_ _ a-. . . . . . . . . . . . . . . .

AI3Ti(Si )

;.a

Ti

Fig. 21. Partial view of the three-dimensional 1000°C A1-C-Si-Ti isotherm showing the solid-solid phase equilibria in the Si-rich corner.

$i

(at-%)

4020-~

10-

a

AI

SiC~

"~

i(Al'si)2 i'r"

~.~/Ti5Si3Cx TiqSiC2

Al4~3 ~ -

/~

/"° AI 0

0,5

~

T' 1

,5

-

Ti (at-%)

Fig. 19. Characteristic lines of the liquidus surface of the A1-C-SiTi system at 1000°C: conic projection£rom the C apex onto the AI-Ti-Si plane.

been constructed from the above-reported results, is reproduced in Fig. 17. On this general view in true perspective, the pure components A1, C, Si and Ti occupy the four apices of the large equilateral tetrahedron limiting the system. On the six edges joining these apices lie the binary carbides A14C3, SiC, TiC, the binary silicides TiSiz, TiSi, TisSi3 and the aluminides A13Ti, TiA1 and Ti3AI. The ternary compounds Ti2A1C, Ti3A1C, T, = Ti3SiC2 and T2 = Ti(Al,Si)2 lie in the triangular faces of the large tetrahedron, faces which correspond to the four ternary subsystems shown in Figs. 1-4 (for clarity, c¢-Ti and fl-Ti have been omitted). The different line segments drawn through the tetrahedron represent the tie lines joining conjugate phases. Conventionally, thick clear lines represent

J.C. Viala et aL / Materials Science and Engineering A229

(1997) 9 5 - i 1 3

109

solid-solid equilibria and dark thin lines solid-liquid equilibria, the liquid lying along the AI-Si edge from pure A1 to a point corresponding to 44 at.% Si.

:

6.1.1. Solid-liquid equilibria In Fig. 18, we have reproduced the part of the 1000°C A I - C - S i - T i isotherm showing the spatial arrangement of the different solid-liquid phase fields. Eight solid phases are in equilibrium with the liquid: these are elemental silicon and the seven compounds A13Ti, SiC, TiC, A14C3, Ti3SiC2, TisSi3Cx and Ti(A1,Si)2. Up to now, it has been convenient to consider the liquid as an AI-Si alloy containing 0-44 at.% Si. In fact, this liquid is a quaternary alloy that can also contain up to about 1.5 at.% Ti (according to EPMA analyses) and in the order of 30 at. pprn C [10]. Owing to the very low solubility of carbon in this quaternary alloy, it was possible to construct the projection, from the C apex onto the AI-Si-Ti plane, of the boundary lines running on the surface limiting the liquid in the 1000°C A I - C - S i - T i isotherm.

._...._/

}r

Ti3SiC2 / :i

_i~

Ti~Si3Cx "

'iSi/

Fig. 24. Projection from the C apex onto the AI-Si-Ti triangle of the solidus and solvus of the ternary compound Ti(A1,Si)2.

C~

\\

Ti3A'~

." . x

Fig. 25. Projection from the A1 apex onto the C - S i - T i triangle of the solidus and solvus of the phase Ti3SiC2.

AlaTi(Si)

TiA1

Ti3AI a-Ti 13-Ti

Fig. 22. Partial view of the three-dimensional 1000°C A 1 - C - S i - T i isotherm showing the solid-solid phase equilibria in the Ti-rich corner.

Fig. 23. Projection from the C apex onto the A1-Si-Ti triangle of the solidus and solvus of the phase A13Ti.

Such a projection, schemed in Fig. 19, clearly shows the liquidus surfaces of the eight above-mentioned solid phases. These liquidus surfaces correspond to twophase equilibria which are bivariant under a constant pressure (1 atm) and at a given temperature (1000°C). This means that these equilibria are not modified when the composition of the liquid moves within the limits of the liquidus surface of each conjugate solid (correspondingly, the composition of the solid can also move on its solidus surface). Two different kinds of bivariant solid-liquid equilibria can be distinguished. Among the first kind are the phase equilibria observed between the liquid and a solid phase located in one of the three ternary subsystems having A1 as an apex, i.e., Si, A13Ti, SiC, TiC, AI4C 3 o r Ti(A1,Si)2. Such two-phase equilibria naturally derive by extension from the A 1 - C - S i - T i quaternary system of phase equilibria that already exist in the ternary subsystems. For that reason, these phase equilibria could easily be predicted a priori, although the case of titanium carbide appears rather singular. In effect, this carbide, which is in equilibrium at 1000°C with liquid aluminum in the A I - C - T i ternary subsys-

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or.C. Viala et aL / Materials Science and Engineering A229 (t997) 95-113

tem, remains in equilibrium in the A I - C - S i - T i quaternary system with an Al-based liquid that can contain up to 11.5 at.% silicon. Of the second kind are the phase equilibria between the liquid and one of the two phases Ti3SiC2 and TisSi3C~ belonging to the C - S i - T i ternary subsystem. In contrast to the previous ones, these phase equilibria of the second kind could not be predicted a priori since they do not derive from any known equilibrium. In this regard, it is worth noting that the ternary compound Ti3SiC2 possesses a wide liquidus: in effect, this compound is in equilibrium with an N-based liquid containing 1-19 at.% silicon. As for TisSi3Cx, it is in equilibrium only with a liquid containing 8-10.5 at.% silicon, but the fact that this phase possesses a liquidus is itself very surprising. Effectively, the silicide TisSi3 from which TisSi3C~ derives is not in equilibrium with any liquid in the Al-Si-Ti ternary subsystem. This suggests that dissolution of carbon in the binary silicide TisSi3 may strongly stabilize this phase.

Fig: 26. Projection from the Si apex onto the A1-C-Ti triangle of the solxqas of the ternary carbide Ti2AIC.

TiC

Ti58i3C x

Ti3A I

Fig. 27. Projection from the Si apex onto the A1-C-Ti triangle of the sotvus of the ternary carbide Ti3A1C.

Along the boundary lines drawn in Fig. 19, the liquid is in equilibrium with two conjugate solid phases, giving rise to 15 different three-phase equilibria. At 1000°C and under 1 atm, these three-phase liquid-solid equilibria are monovariant. In the same way as previously, they can be classified in two groups. A first group consists of six phase equilibria that derive naturally from equilibria already existing in the three ternary subsystems having A1 as an apex. Equilibria such as TiC-Ai3Ti-L(0_I~ or TiC-AI4C3-L(o_Hs) belong to this first group. The second group consists of nine equilibria that have no equivalent in any ternary subsystem limiting the A I - C - S i - T i tetrahedron. Among these nine equilibria, one can mention as examples SiC-Ti3SiC2-L(la_~9), SiC- Ti(A1,Si)2-L(ag_44) or TisSi3Cx-Ti3SiCa-L(8_10.5~. Concerning this second group, it may be noted that A14C3 and Ti3SiCa coexist with a liquid containing 11.5-12 at.% Si: as a consequence, TiC and SiC cannot be simultaneously in equilibrium with any liquid, although these two carbides are in equilibrium in the solid state. Similarly, TiC and TisSi3Cx, which are in equilibrium in the solid state, cannot be simultaneously conjugate with any liquid, due to the existence of the three-phase equilibrium Ai3Ti-Ti3SiC2-L(I_~. In the eight points where three boundary lines intersect (Fig. 19), a liquid with a particular composition is simultaneously in equilibrium with three conjugate solid phases. Such four-phase equilibria are invariant at 1000°C under 1 atm. Consequently, in the three-dimensional representation given in Fig. t8, each of these points of the Al-Si line is the apex of a tetrahedron, the three other apices of which correspond to a particular composition in the homogeneity range of three conjugate solids. These tetrahedrons are TiC-A13TiTi3SiCz-L(I~, TisSi3CzAi3Ti-Ti3SiC2-L(8~, TisSi3CxA13Ti-Ti(Al,Si)2- L(lo~, TisSi3CzTi3SiC2-Ti(A1,Si)2L(lo.:), TiC-A14C3-Ti3SiC2-L(::.5~, SiC-A14C3Ti3SiC2-L(la~, SiC-Ti3SiC2-Ti(AI,Si)2-L(~9~ and SiCSi-Ti(A1,Si)2-L(44~. It can be noticed that, due to the fact that the phase Ti3SiC2 is in equilibrium with a wide range of liquid compositions (from 1 to 19 at.% Si), this phase is present in six of the eight four-phase liquidsolid equilibria. Finally, it can be remarked that there is no liquid simultaneously in equilibrium with TiC, Al3Ti and TisSi3C~, whereas a three-phase equilibrium exists between these compounds in the solid state. The same remark can be made for the conjugate solid phases TiC, Ti3SiC2 and TisSi3C~. 6.1.2. Solid-solid equilibria In order to facilitate the representation of the different phase equilibria that tend to establish between solid phases in the A I - C - S i - T i system at 1000°C under 1 atm, we have reproduced in Figs. 20-22 three partial views of the tetrahedron presented in Fig. 17.

J.C. Viala et al./ Materials Science and Engineering A229 (1997) 9 5 - i i 3

111

~lO

AI

AI3Ti

'l'iAI

Ti3AI

c~Ti

~Ti

Ti

Fig. 28. The 1000°CA1-Ti-SiC isopleth. Fig. 20 represents the solid-solid phase equilibria existing in the carbon-rich corner of the A I - C - S i - T i system. Two invariant four-phase equilibria are remarkable: they correspond to two tetrahedrons having A14C3, SiC, a carbon-rich composition of TiC as base triangle and either C or Ti3SiC2 as apex. It is recalled that a monovariant three-phase equilibrium such as A14C3-SiC-TiC should strictly be represented by a volume limited by two triangles. In the present case, however, the variations in composition of the three phases involved within the equilibrium volume are small and this volume is reduced to a simple triangle. In Fig. 21 are shown the solid-solid phase equilibria in the silicon-rich corner of the A I - C - S i - T i system. In the foreground of this figure are represented two different three-phase monovariant equilibria: TiC-A13TiTisSi3Cx, which is reduced to a simple triangle, and TisSi3Cx-A13Ti-TiSi, which appears as a volume leaning on a line of the solvus of TisSi3C~. Four tetrahedrons can be observed at middle distance, representing the four four-phase equilibria: TiC-A13Ti-Ti3SiC 2TisSi3Cx, TisSi3Cx-Ti(A1,Si)2-TiSi-A13Ti, TisSi3CxTiSi-Ti(A1,Si)2-TiSi2 and TisSi3Cx-Ti3SiC2Ti(A1,Si)2-TiSi2. In the background, two tetrahedrons represent the two four-phase equilibria SiC-Ti3SiC2Ti(A1,Si)2-TiSi2 and SiC-Si-Ti(A1,Si)2-TiSi2. Fig. 22 illustrates the inter-relation of the different phase regions in the part of the A I - C - S i - T i tetrahedron located between the TiC-A13Ti-TisSi3C~ triangle and the Ti apex. This partial section is consistent with

the semischematic diagram constructed at 1200°C by Abboud and West [33]. It is worth noting that at 1000°C under 1 atm, all the phases lying in the front triangle TiC-A13Ti-Ti are in equilibrium with different compositions of the line compound TisSi3Cx. This remains true (law of adjoining phase regions) for the c~-Ti and fl-Ti phases which have been included in Fig. 22.

6.1.3. Solidus and solvus projections By analogy with the liquidus projection given in Fig. 19, the characteristic lines of the solidus and solvus of some compounds have been projected from an apex onto the appropriate opposite triangle of the A I - C Si-Ti tetrahedron. On such projections presented in Figs. 23-2, a surface corresponds to a bivariant twophase equilibrium, a line separating two surfaces to a monovariant three-phase equilibrium and a point in which three lines intercept to an invariant four-phase equilibrium. Although schematic, such projections offer the advantage of providing a two-dimensional representation of the main phase equilibria that exist in the A 1 - C - S i - T i quaternary system at 1000°C under 1 atm. More especially, it can easily be seen on these projections how the solid-solid and solid-liquid equilibria are inter-related.

6.1.4. Restrictive conditions At this stage of the discussion, it will be retailed that simplifications have been made to clarify the represen-

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J.C. Viala et al./ Materials Science and Engineering A229 (1997) 95-113

tation of the A I - C - S i - T i isobaric-isothermal section reported in Fig. 17: the homogeneity ranges of the phases considered have been neglected, except for A13Ti , TiA1, Ti3A1, TiC and TisSi3Cx that have been schematically represented as line compounds, the phases A12Ti, A15Ti2, Ti3Si and TisSi 4 have been omitted (and also a-Ti and/3-Ti, but the case of these solid solutions has already been discussed), and the phase equilibria taken into consideration in the A1-C-Si ternary subsystem are only metastable. Consequently, the limits of numerous phase regions are not very accurately defined and, due to omitted phases, certain phase regions could divide into sub/'egions. With these reservations, the tetrahedron we have constructed and the two-dimensional sections or projections that have been derived from it remain powerful models for describing the thermodynamic aspect of most of the reactions likely to occur in the A I - C - S i - T i system at 1000°C under 1 atm. This is particularly the case for the chemical interactions between silicon carbide (SIC) and the A1-Ti alloys that will now be discussed in more detail.

6.2. Thermodynamic aspect of the chemical interaction between SiC and Al-Ti alloys The three isopleths A13Ti-Si-C , Ti3AI-Si-C and TiAI-Si-C reported in Figs. 6, 15 and 16 account for all the phase equilibria that tend to be established when one of the stoichiometric titanium aluminide A13Ti, Ti3A1 or TiA1 reacts at 1000°C under 1 atm with pure SiC or with SiC + Si or SiC + C mixtures. In addition to these data, a fourth isopleth having A1, Ti and SiC as apices has been constructed. This new two-dimensional section, shown in Fig. 28, represents all the phase equilibria that are to be attained after complete reaction at 1000°C between pure SiC and any A1-Ti alloy, whatever the composition of the alloy and the atomic fraction of SiC. It is the latter isopleth that will be used to illustrate the main thermodynamic principles governing the high-temperature chemical interaction between SiC and AI-Ti alloys. Firstly, it can be remarked that, contrary to the A I - C - N i - S i system where equilibrium conditions exist at high temperature between SiC and the nickel aluminide A1Ni [34], there is no phase region in which SiC and an AI-Ti alloy coexist. This means that at 1000°C, reactions are likely to proceed in any SiC:AI-Ti alloy mixture until one at least of the two constituents is completely consumed. If the initial SiC:A1-Ti alloy mixture contains sufficient amounts of SiC (i.e., from more than about 20 at.% SiC for Al-rich alloys to more than about 50 at.% SiC for Ti-rich alloys), an unreacted fraction of this carbide will be recovered after complete reaction. This nnreacted SiC will be in equilibrium either with solid

phases if the A1 content in the starting A1-Ti alloy is less than about 10 at.% or with solid phases and an AI-Si liquid if the starting AI-Ti alloy contains more than about 10 at.% AI. According to Fig. 28, the solid-solid four-phase region corresponding to the former case is SiC-Ti3SiC2-Ti(A1,Si)z-TiSi2 and the solid-liquid phase regions corresponding to the latter cases are SiC-Ti3SiC2-Ti(A1,Si)z-L(19) when the A1 content in the initial AI-Ti alloy ranges from 10 to about 57 at.%, SiC-Ti3SiCz-L(lz_19) when this A1 content ranges from 57 to about 70 at.%, and SiCTi3SiCz-A1,C3-L(12) for A1 contents higher than about 70 at.%. These conclusions are in agreement with the results reported at 1000°C for SiC crystals joined with foils of reactive metals [35]. If the initial SiC:AI-Ti alloy mixture is SiC-poor (from less than about 20 at.% SiC for Al-rich alloys to less than about 50 at.% SiC for Ti-rich alloys), this carbide will be entirely decomposed and a great number of phase equilibria are to be attained after complete reaction, as shown by Fig. 28. As long as solid-state reactions are concerned, silicon and carbon resulting from the decomposition of SiC will be converted into the silicide TisSi3Cx and into one of the binary or ternary carbides TiC, Ti3A1C or TiaA1C. TiC will be formed with the Ti-rich AI-Ti alloys and Ti2A1C with the M-rich AI-Ti alloys (Ti3A1C constitutes an intermediate case). It is important to note that formation of these compounds (silicide and carbides) consumes much more titanium than aluminum: consequently, the remaining AI-Ti alloy will always be enriched in aluminum. This enrichment may result in the formation of an Al-based liquid and, in that case, the complete decomposition of SiC by this liquid produces TiC and Ti3SiC2 or A14Cy

7. Conclusion

This work was undertaken with the aim of obtaining a better understanding of the reaction processes likely to develop at high temperature at the interface between silicon carbide reinforcements and titanium aluminide matrices. With the experimental determination of the phase equilibria in the A I - C - S i - T i quaternary system at 1000°C under a pressure of 1 atm, the thermodynamic grounds of this interface chemistry have been established. Further studies are now in progress on SiC/titanium aluminide finite and semi-infinite diffusion couples to characterize the interfacial zones formed by reaction in out-of-equilibrium conditions, to relate the reaction layer sequences experimentally observed with the A I - C - S i - T i phase diagram using the diffusion path concept and to propose reaction mechanisms describing the growth of these zones in terms of multiphase volume diffusion.

J.C. Viala et al./ Materials Science and Engineering A229 (1997) 95-I13

Acknowledgements Support from the French Ministry for Research and Technology is gratefully acknowledged. Characterization by SEM and EPMA was achieved using the equipment of CMEABG, Universit~ Lyon 1.

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